Porous Anodic Aluminum Oxide: Anodization and Templated

Review pubs.acs.org/CR

Porous Anodic Aluminum Oxide: Anodization and Templated Synthesis of Functional Nanostructures Woo Lee*,†,‡ and Sang-Joon Park† †

Korea Research Institute of Standards and Science (KRISS), Yuseong, 305-340 Daejeon, Korea Department of Nano Science, University of Science and Technology (UST), Yuseong, 305-333 Daejeon, Korea

‡

6.2.3. Morphological Instability 6.3. Steady-State Pore Formation 6.3.1. Joule’s Heat-Induced Chemical Dissolution 6.3.2. Field-Assisted Oxide Dissolution 6.3.3. Average Field Model for Steady-State Pore Structure 6.3.4. Direct Cation Ejection Mechanism 6.3.5. Flow Model for Steady-State Pore Formation 7. Self-Ordered Porous Anodic Aluminum Oxide (AAO) 7.1. Mild Anodization (MA) 7.2. Hard Anodization (HA) 7.3. Pulse Anodization (PA) 7.4. Cyclic Anodization (CA) 7.5. Anodization of Thin Aluminum Films Deposited on Substrates 8. Long-Range Ordered Porous AAO 9. AAO Template-Based Synthesis of Functional Nanostructures 9.1. Electrochemical Deposition (ECD) 9.2. Electroless Deposition (ELD) 9.3. Sol−Gel Deposition 9.4. Surface Modification 9.5. Template Wetting 9.6. Mask Techniques 9.7. Chemical Vapor Deposition (CVD) 9.8. Atomic Layer Deposition (ALD) 10. Closing Remarks and Outlook Author Information Corresponding Author Notes Biographies Acknowledgments Abbreviations References

CONTENTS 1. Introduction 2. Types of Anodic Aluminum Oxide (AAO) 3. Ionic Conduction in Anodic Oxide Films 3.1. High-Field Conduction Theory 3.2. Elementary Interfacial Reactions 3.3. Transport Numbers 3.4. Stress-Driven Ionic Transport 4. Electrolytic Breakdown 4.1. Factors Influencing Breakdown 4.1.1. The Nature of Anodized Metal 4.1.2. Electrolyte Conditions 4.1.3. Current Density (j) 4.1.4. Other Factors Influencing Breakdown 4.2. Models for Breakdown 4.2.1. Electron Avalanche Multiplication 4.2.2. Stress-Driven Breakdown 5. Structure of Porous Anodic Aluminum Oxide (AAO) 5.1. General Structure 5.1.1. Pore Diameter (Dp) 5.1.2. Interpore Distance (Dint) 5.1.3. Barrier Layer Thickness (tb) 5.2. Structure of Pore Wall (Anion Incorporation) 5.3. Effect of Heat Treatments 6. Growth of Porous Anodic Aluminum Oxide (AAO) 6.1. Stress Generation in Anodic Oxide Films 6.1.1. Volume Expansion 6.1.2. Stress Measurements 6.1.3. Effects of External Stresses on Pore Growth 6.2. Initial-Stage Pore Formation 6.2.1. Qualitative Description on Pore Formation 6.2.2. Kinetics of Porosity Initiation © 2014 American Chemical Society

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1. INTRODUCTION In ambient atmospheres, aluminum becomes rapidly coated with a compact 2−3 nm thick oxide layer. This native oxide layer prevents the metal surface from further oxidation. Because of the surface native oxide, aluminum generally has good corrosion resistance. However, local corrosion of metal can occur in rather aggressive outdoor environments, containing

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Received: January 2, 2014 Published: June 13, 2014 7487

dx.doi.org/10.1021/cr500002z | Chem. Rev. 2014, 114, 7487−7556

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discussed (section 5). Anodization of aluminum is a volume expansion process, and thus is accompanied by mechanical stresses. Recent studies have indicated that the stresses have profound implications not only on the ionic transport, but also on the self-ordering behavior of oxide nanopores. We will discuss in detail the effect of stress on pore growth (section 6.1), the kinetics of pore initiation, and morphological instability associated with the early stage of anodization (section 6.2), and recent models describing steady-state pore formation (section 6.3). After that, recent progress on anodization of aluminum used in fabricating self-ordered porous AAO and also for engineering internal pore structures will be discussed (section 7). In addition, various approaches to long-range order porous AAO will be reviewed (section 8). In the last part of this Review (section 9), various chemical approaches for the syntheses of low-dimensional functional nanostructures and the fabrications of advanced nanodevices will be discussed. These approaches include electrochemical deposition (ECD), electroless deposition (ELD), sol−gel deposition, surface modification, template wetting, shadow mask techniques, chemical vapor deposition (CVD), and atomic layer deposition (ALD). Chemistry issues encountered in the template-based synthesis of functional nanostructures will be discussed in detail. Finally, we will present the challenges and future prospects of the field (section 10).

corrosive chemicals (e.g., chlorides or sulfates). In 1857, Buff first found that aluminum can be electrochemically oxidized in an aqueous solution to form an oxide layer that is thicker than the native one.1 This phenomenon has been called “anodization” because the aluminum part to be processed constitutes the anode in an electrolytic cell. In the early 1920s, the phenomenon observed by Buff was exploited for industrial scale applications, for example, protection of seaplane parts from corrosive seawater.2 In general, the anodic aluminum oxide (AAO) films form with two different morphologies (i.e., nonporous barrier-type oxide films and porous-type oxide films) depending mainly on the nature of the anodizing electrolyte.3 Because the process was first implemented for protection purposes, the anodization of aluminum and its alloys, particularly porous-type anodization, has received considerable attention in the industry because of its extensive practical applications. Many desirable engineering properties such as excellent hardness, corrosion, and abrasion resistance can be obtained by anodizing aluminum metals in acid electrolytes.4 In addition, due to its high porosity, the porous oxide films formed on the metals serve as a good adhesion base for electroplating, painting, and semi-permanent decorative coloration. The anodized products can be easily found in electronic gadgets, electrolytic capacitors, cookware, outdoor products, plasma equipment, vehicles, architectural materials, machine parts, etc. Recently, this nearly century-old industrial process has been drawing increasing attention from scientists in the field of nanotechnology. This trend originated with the seminal works of Masuda and co-workers, who reported on self-ordered porous AAO in 19955 and the subsequent development of the two-step anodization process in 1996.6 Porous AAO film grown on aluminum is composed of a thin barrier oxide layer in conformal contact with aluminum, and an overlying, relatively thick, porous oxide film containing mutually parallel nanopores extending from the barrier oxide layer to the film surface.7 Each cylindrical nanopore and its surrounding oxide region constitute a hexagonal cell aligned normal to the metal surface. Under specific electrochemical conditions, the oxide cells self-organize into hexagonal closepacked arrangement, forming a honeycomb-like structure.5−7 Pore diameter and density of self-ordered porous AAOs are tunable in wide ranges by properly choosing anodization conditions: pore diameter = 10−400 nm and pore density = 108−1010 pores cm−2. The novel and tunable structural features of porous AAOs have been intensively exploited for synthesizing a diverse range of nanostructured materials in the forms of nanodots, nanowires, and nanotubes, and also for developing functional nanodevices. The objective of this Review is to provide a solid information source for researchers entering this field and to establish a broad and deep knowledge base. This Review introduces the fundamental electrochemical processes associated with anodic oxidation of aluminum, and discusses the recent progress on anodization of aluminum for the development of ordered porous AAOs, and nanotechnology applications of porous AAOs. We organize this Review as follows: after discussing the growth characteristics of two different types of AAOs (section 2), we will describe the theory of ionic conductions and elementary interfacial reactions (section 3), followed by electrolytic breakdown (section 4) to understand the fundamental electrochemistry associated with anodic oxidation of aluminum. Next, the electrochemical factors defining the geometric and chemical structures of porous AAOs will be

2. TYPES OF ANODIC ALUMINUM OXIDE (AAO) Anodization of aluminum in aqueous electrolytes forms anodic oxide films with two different morphologies, that is, the nonporous barrier-type oxide films and the porous-type oxide films. The chemical nature of the electrolytes mainly determines the morphology of AAOs.3,7,8 A compact nonporous barrier-type AAO films can be formed in neutral electrolytes (pH 5−7), such as borate, oxalate, citrate, phosphate, adipate, tungstate solution, etc., in which the anodic oxide is practically insoluble.9,10 Meanwhile, porous-type AAOs are formed in acidic electrolytes, such as selenic,11 sulfuric,12 oxalic,12 phosphoric,7,12,13 chromic,12,14 malonic,12,15−17 tartaric,12,18 citric,12,17−20 malic acid,12,18 etc., in which anodic oxide is slightly soluble. Early models describing anodic oxide growth were developed on the basis of the barrier-type oxide.21−24 Moreover, in the early stage of porous-type oxide growth, the formation of the initial barrier oxide is followed by the emergence of incipient pores. Therefore, in this Review, we will mention the barrier-type oxide growth to the extent needed for understanding porous-type oxide formation. Some excellent review articles covering the barrier-type anodic oxide films are given in refs 3 and 25. The two types of anodic oxides (i.e., barrier- vs porous-type AAO) differ in their oxide growth kinetics. In the case of barrier-type oxide formation under potentiostatic conditions (i.e., U = constant), current density (j) decreases exponentially with time (t). Correspondingly, the film growth rate decreases almost exponentially with time (t), which places a limit on the maximum film thickness obtainable for barrier-type AAO films (Figure 1). It has been experimentally verified that the thickness of barrier-type film is directly proportional to the applied potential (U). On the other hand, current density (j) in porous-type anodization under potentiostatic conditions remains almost constant within a certain range of values during the anodization process, due to the constant thickness of the barrier layer at the pore bottom. The thickness of the resulting porous oxide film is linearly proportional to the total amount of 7488


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aluminum, new oxide forms above and below the marker layer (Figure 2b). The marker layer is located at a depth of 40% of the film thickness in a plane corresponding to that of the original metal surface. On the other hand, when a barrier-type film is formed at 60% current efficiency (ηj), the plane of the marker layer is immobile and 40% of the Al3+ cations are shed into the electrolyte via direct cation ejection mechanism without contributing to the oxide formation (Figure 2c). In this case, anodic oxide grows at the metal/oxide interface via the inward migration of O2−/OH− ions. When porous-type anodic oxide forms at 60% current efficiency, the marker plane is located above that of the original metal surface (Figure 2d). In this case, the metal/oxide interface is also the oxide growth front, and 40% of Al3+ ions are ejected into the solution. Because cations are being shed into the electrolyte, the current efficiency (ηj) of porous-type oxide growth is typically much lower than that of the barrier-type. Accordingly, the Pilling−Bedworth ratio (PBR = the ratio of molar volume of the grown oxide to molar volume of the consumed metal; see section 6.1.1) for the initial barrier oxide formation in poroustype oxide growth at the early stage of anodization is lower than that for barrier-type oxide growth: PBR = 0.90 for porous-type oxide growth at ηj = 53.5% in phosphoric acid solution and PBR = 1.7 for barrier-type oxide growth at ηj = 100% in neutral adipate solution.9,33 Shimizu et al.33 suggested that the initial barrier oxide grows under increasing tensile stress (PBR < 1), which causes local oxide cracking most probably at the randomly present metal protrusions. The generated surface cracks were considered to be local paths for electrolyte penetration, causing non-uniform local thickening of the initial barrier oxide. Non-uniform thickening of the initial oxide causes concentration and redistribution of the current lines into the relatively thin oxide regions between the protrusions (i.e., a local increase in electric field, E). Consequently, localized scalloping of the metal/oxide interface takes place. Shimizu et al.33 pointed out that the non-uniform thickening of anodic oxide (i.e., morphological instability) in the initial barrier oxide is “one of the most distinctly different growth features between porous- and barrier-type AAO films”. Unlike porous-type oxide growth in acid electrolytes, anodic oxides in neutral electrolytes grow highly uniformly on surface finished aluminum, maintaining flat metal/oxide and oxide/electrolyte interfaces. Even the smoothing of initially rough aluminum surfaces during the growth of barrier oxide films has been experimentally observed.34

Figure 1. Two different types of anodic aluminum oxide (AAO) formed by (a) barrier-type and (b) porous-type anodizations, along with the respective current (j)−time (t) transients under potentiostatic conditions.

charge (i.e., anodization time, t) involved in the electrochemical reaction. Radiotracer studies, employing an immobile marker (125Xe), have indicated that, in the case of barrier-type oxide formation, anodic alumina grows simultaneously at the oxide/electrolyte interface and at the metal/oxide interface, through Al3+ egress and O2−/OH− ingress, respectively, under a high electric field (E).26−28 In the case of porous-type anodic alumina formation, on the other hand, oxide grows at the metal/oxide interface via the inward migration of O2−/OH− ions. 18O tracer studies have shown that outwardly migrating Al3+ cations do not contribute to the oxide growth at the oxide/electrolyte interface, but are all shed into the anodizing electrolyte via direct ejection mechanism (see section 6.3.4).10,29−31 Otherwise, egressing Al3+ ions would form anodic alumina at the oxide/electrolyte interface to heal any developing or embryonic pores there. Schematic diagrams illustrating the dimensional changes of aluminum during the barrier-type and porous-type anodic oxide formation are shown in Figure 2.32 An immobile marker layer is implanted into the starting aluminum with a native oxide layer (Figure 2a). When a barrier-type film is formed at 100% current efficiency (ηj) by anodization of the marker-implanted

Figure 2. Schematic diagrams illustrating dimensional changes of an aluminum specimen following anodizing. (a) Initial aluminum with a thin airformed oxide film. The red dashed line represents an immobile marker layer implanted into the initial aluminum with a thin air-formed oxide film. (b) Anodized at 100% efficiency with formation of a barrier-type anodic film. (c) Anodized at just above 60% efficiency with formation of a barriertype anodic film. (d) Anodized at 60% efficiency with formation of a porous anodic film. Reproduced with permission from ref 32. Copyright 2006 The Electrochemical Society. 7489


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in the oxide is high enough (e.g., 106−107 V cm−1), the ionic current density (j) can be expressed as25

3. IONIC CONDUCTION IN ANODIC OXIDE FILMS 3.1. High-Field Conduction Theory

⎛ W ⎞⎛ αazFE ⎞ ⎟⎜ ⎟ j = vρa exp⎜ − ⎝ RT ⎠⎝ RT ⎠

When a valve-metal is anodized under either potentiostatic or galvanostatic condition, anodic oxide film forms on the metal. For anodizing aluminum (Al) and tantalum (Ta), an empirical exponential dependence of the ionic current density (j) on the electric field (E) is established. Ionic current density (j) under high-field conditions, which is the case for anodic oxide growth, can be associated with the movement of charged ions in the barrier oxide, and can be related to the potential drop (ΔU) across the barrier oxide through the exponential law of Güntherschulze and Betz, as follows:21,35 j = j0 exp(βE) = j0 exp(β ΔU /tb)

(2)

where v is the hopping attempt frequency of the ion, ρ is the density of concentration of mobile charge in C cm−3, a is the hopping inter-distance, W is the hopping activation energy at zero field, α is a parameter describing the asymmetry of the activation barrier at non-zero field, z is the valence of the mobile ions, and F is Faraday’s constant. From eqs 1 and 2, the following relations can be obtained: ⎛ W ⎞ ⎟ j0 = υρa exp⎜ − ⎝ RT ⎠

(1)

where j0 and β are material-dependent constants at a given temperature, and ΔU/tb is the effective electric field (E, typically 106−107 V cm−1) impressed on the barrier layer with thickness tb. For anodic alumina, a large range of j0 and β values has been reported: j0 = 3 × 104 to 1 × 10−18 A cm−2 and β = 0.1 × 10−6 to 5.1 × 10−6 cm V−1.25 For anodic oxidation of metal in an electrolyte, three theories based on the following possible rate-determining steps for oxide formation have been developed:3 ion transfers (i) across the metal/oxide interface (Mott−Cabrera theory),23,24 (ii) through the oxide bulk (Verwey theory),22 and (iii) across the oxide/ electrolyte interface (Dewald theory).36,37 In the point defect model of Macdonald et al.,38 the oxide film is assumed to contain a high concentration of non-interacting positive and negative point defects, and the rate-determining step for the oxide growth is assumed to be the transport of metal and oxide vacancies across the oxide film. All of these theories can explain the empirical exponential relationship proposed by Güntershultz and Betz. On the other hand, transient experiments favorably indicate that the rate-determining step is the movement of charged ions within the oxide.25 On the basis of the rate-determining movement of ions within the oxide, the high-field model relates the parameters j0 and β in eq 1 to the nature of oxide materials. The high-field conduction model is based on a hopping mechanism, in which the activation energy for hopping ions is dependent on electric field E (Figure 3).25 Ions at regular sites or interstitial positions jump to vacancies or other interstitial positions in their neighborhood. The model assumes that the oxide is defect-free and of homogeneous composition. When the electric field (E)

β=

αazF RT

(3) (4)

Because the parameter a can be related to the inter-atomic distance in the oxide, one can expect that the electric field strength (E) increases when the oxygen ion density increases (i.e., a decrease in parameter a) provided that the other parameters are constant. Equation 1 can be modified to obtain a Tafel equation:

ln j = ln j0 + βE

(5)

For a constant oxide thickness tb, a constant Tafel slope β is obtained. The electric field (E) in the oxide can be related to the applied (or measured, in the potentiostatic condition) electrode potential (U). The measurable potential drop between the metal and the electrolyte is equal to U = ΔU + Φm/o + Φo/e

(6)

where ΔU is the potential drop in the oxide, and Φm/o and Φo/e are the potential drops at the metal/oxide and oxide/electrolyte interfaces, respectively.39 In a typical anodization, the potential drops at the metal/oxide and oxide/electrolyte interfaces are quite small, as compared to the several tens of volts of potential drop in the oxide (i.e., ΔU ≫ Φm/o + Φo/e). Therefore, the following approximation for the electric field (E) is possible for the high-field ionic transport: E = ΔU /tb ≈ U /tb

(7)

where tb is the thickness of oxide. 3.2. Elementary Interfacial Reactions

As was already discussed in section 2, it is now widely accepted that for the anodic growth of alumina both Al3+ cations and oxygen-containing anions (e.g., O2− or OH−) are mobile within the anodic oxide under high electric field (E).10,26−28,40 Al3+ ions migrate outwardly toward the oxide/electrolyte interface, while O2− or OH− anions move inwardly toward the metal/ oxide interface. Therefore, one can consider both (i) the metal/ oxide and (ii) the oxide/electrolyte interfaces as the growth front of anodic oxide during anodization of a valve-metal. For anodizing aluminum, the following elementary reactions are considered to be possibly occurring at the interfaces (Figure 4). (i) At the metal/oxide interface: Figure 3. Influence of the electric field strength (E) on the activation energy of hopping ions. Reproduced with permission from ref 25. Copyright 1993 Elsevier. 7490

3+ Al → Al(ox) + 3e−

(8)

3+ 2− 2Al(ox) + 3O(ox) → Al 2O3

(9)


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Figure 4. Schematic diagrams showing elementary interfacial reactions for (a) barrier-type and (b) porous-type anodic oxide.

and t− = 1 − t+ for anion. Transport numbers can be determined by employing a “marker layer”, whose position in the anodic oxide film indicates the extent of oxide that was formed at each interface. If the metal ions are the only mobile species, new oxide should be formed at the oxide/electrolyte interface on top of the marker layer. On the other hand, if oxygen anions are the only mobile species, the new oxide should be formed at the metal/oxide interface below the marker layer. Davies et al.26 stated that the ideal marker atoms for determination of transport numbers should fulfill the following requirements: “The markers should be (i) uncharged, so that they do not migrate in the oxide under the influence of the applied field; (ii) large in size, so that they do not diffuse significantly within the oxide lattice; (iii) present in trace amount, so that the macroscopic properties of the tagged oxide remain unaltered; and (iv) detectable, in order to assess their depth in the oxide.” These conditions can be satisfied by implanting radioisotopes 125 Xe inert gas atoms or 222Rn, which are heavier than typical valve-metals and oxygen, in a preformed thin oxide film and subsequently anodized.26,27,42 Radioactive tracers allow the position of the buried marker to be assessed by monitoring the energy of emitted α- or β-particles.26,27,42,43 Other techniques to measure the buried marker position in oxide include Rutherford backscattering spectrometry (RBS)40 or direct observation of voids formed by implanted Xe by employing cross-sectional transmission electron microscopy (TEM).28,44 A representative cross-sectional TEM image showing an immobile Xe marker layer is given in Figure 5. The sample in the figure was formed in near-neutral potassium phosphate electrolyte at a high current efficiency (ηj ≈ 100%).44 The approximately 10nm-thick straight Xe marker layer is located at about the midpoint of the film. The anodic oxide above the marker layer formed by the field-driven outward migration of Al3+ ions and that beneath the marker layer by the field-driven inward migration of oxygen carrying anions, O2−/OH−. Assuming that all egressing Al3+ ions contribute to the oxide formation, the cation transport number was directly estimated to be t+ = 0.49.44 For anodizing conditions under which oxide grows with appreciable metal dissolution, however, TEM-based direct measurement may underestimate the cation transport number. In such cases, Al3+ ions dissolved in anodizing electrolyte should be quantified to estimate the equivalent oxide thickness. Davies et al.26 pointed out that the location of an immobile marker in anodic oxide markedly depends on the anodization conditions, such as current density (j) and the nature of

(ii) At the oxide/electrolyte interface: 3+ 2− 2Al(ox) + 3O(ox) → Al 2O3

(10)

+ 3+ Al 2O3 + 6H(aq) → 2Al(aq) + 3H 2O(1)

(11)

3+ 3+ Al(ox) → Al(aq)

(12)

2− 2O(ox) + O2(g) + 4e−

(13)

− + 2− 2H 2O(1) + O(ox) + OH(ox) + 3H(aq)

(14)

Reactions 9 and 10 correspond to the formation of anodic oxide at the metal/oxide and oxide/electrolyte interfaces, respectively. Reaction 11 describes dissolution of anodic alumina by Joule’s heat-induced oxide dissolution and/or field-induced oxide dissolution, which will be discussed in section 6.3.1 and section 6.3.2, respectively. On the other hand, reaction 12 occurs through field-assisted direct ejection of Al3+ ions from the metal/oxide interface through oxide into the electrolyte, which will be discussed in detail in section 6.3.4. Reactions 11−13 decrease the net current efficiency (ηj) associated with the anodic oxide formation. Reaction 14 describes the heterolytic dissociation of water molecules at the oxide/electrolyte interface, which supplies oxygen anions to the metal/oxide interface to form anodic oxide. By assuming that all oxide anions from the dissolution of Al2O3 at the oxide/ electrolyte interface migrate to the metal/oxide interface to reform Al2O3, and that all oxide anions from the dissociation of water contribute to the oxide formation, Su et al. proposed the following overall reaction at the oxide/electrolyte interface:41 3+ 2− Al 2O3 + nH 2O(1) → 2Al(aq) + (3 − n − x)O(ox) − + + xOH(ox) + (2n − x)3H(aq)

(15)

where n denotes the amount of water dissociated per mole of Al2O3 that is dissolved at the same time. Su et al. claimed the field-dependent nature of the heterolytic dissociation of water in reaction 14 and related the dissociation rate of water to the porosity (P) of AAO, which will be touched upon in section 7.1. 3.3. Transport Numbers

As mentioned in previous sections, anodic oxide formation can occur at both the metal/oxide and the oxide/metal interfaces. The relative amount of mobile ions transported to the oxide forming interfaces is called the “transport number”: t+ for cation 7491


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Figure 5. Cross-sectional transmission electron microscopy (TEM) image showing immobile 125Xe marker layer. The sample (i.e., barriertype anodic oxide film) was formed at a constant current density of 1 mA cm−2 to 100 V in near-neutral potassium phosphate electrolyte. Adapted from ref 44 with permission. Copyright 1987 Taylor & Francis (www.tandfonline.com).

Figure 6. Schematics showing (a) the movement of Al3+ and O2− ions during the re-anodizing process and (b) the corresponding cell potential (U)−time (t) curve. Reproduced with permission from ref 46. Copyright 1978 Elsevier.

electrolyte. The authors performed transport number experiments using radioisotope 125Xe marker atoms and also quantitative analyses on dissolving Al3+ ions during anodization in two different near-neutral electrolytes (i.e., sodium tetraborate and ammonium citrate). According to their experiments, both metal and oxygen ions are mobile during oxide growth. In the borate solution, anodic alumina grew with high current efficiencies (i.e., negligible cation loss), and thus the immobile 125Xe markers were completely buried in the oxide. The mean cation transport number (t+) was estimated to be t+ = 0.58 for anodic alumina formed at the current density range of 0.1−10 mA cm−2. On the other hand, in citrate solution, the amount of aluminum passed into the solution was as high as 40% of the total oxidized metal at low current density, but decreased as the current density increased. The 125 Xe markers in anodic alumina remained very close to the outer surface. The cation transport number varied with current density, from about t+ = 0.37 at 0.1 mA cm−2 to t+ = 0.72 at 10 mA cm−2. In the case of porous AAO, the transport numbers of mobile ions can be estimated by the so-called “pore-filling method”, which was originally used to determine the porosity (P) of porous AAO by Dekker and Diddelhoek.45,46 In this method, aluminum is first anodized in an acid electrolyte to form porous-type anodic oxide and subsequently re-anodized in a neutral electrolyte electrolyte to form barrier-type oxide under a galvanostatic condition. During anodizations, potential (U)− time (t) transients are monitored. During the barrier-type anodizing (i.e., re-anodizing process), new oxide gradually forms simultaneously within the pores and underneath the barrier layer of the pre-formed porous anodic oxide, because both Al3+ and O2− ions contribute to the oxide formation at the metal/oxide and oxide/electrolyte interfaces, respectively. As a result, the cell potential gradually increases with time during the re-anodizing process. Figure 6 schematically shows (a) the movement of Al3+ and O2− and (b) the cell potential (U)−time (t) profile during the re-anodizing process.46 The non-zero value of U at t = 0 is due to the original barrier layer of preformed porous AAO. The complete filling of pores is accompanied by the change of the slope in the U−t curve at time tp due to the sudden increase of the oxide/electrolyte interfacial area. For the time t < tp, the following relation can be obtained:46

⎡ dh+ jM dh− ⎤ ρ ⎢P + ⎥= ⎣ dt dt ⎦ nFk

(16)

−3

where ρ (=2.95 g cm ) is the density of oxide, P is the porosity of porous AAO, dh+/dt and dh‑/dt are, respectively, the rates of the increase of the barrier oxide thickness at the oxide/ electrolyte and metal/oxide interfaces, j is the current density, M is the atomic weight of Al, n (=3) is the number of electrons involved in oxidation reaction, F is Faraday’s constant, and k (=0.505) is the weight fraction of aluminum in the oxide. The cation transport number is given by the ratio of the weight of new oxide formed within the pores per unit time to the total weight of new oxide formed per unit time:46 ⎛ dh+ ⎞ ⎛ dh+ dh− ⎞ t + = ⎜P + ⎟ / ⎜P ⎟ ⎝ dt ⎠ ⎝ dt dt ⎠

(17)

The slopes m1 and m2 of the U−t transient in Figure 6b are given by46 m1 =

1 ⎛ dh+ dh− ⎞ + ⎜ ⎟ AR ⎝ dt dt ⎠

(18)

m2 =

1 ⎛ dh+ dh− ⎞ + ⎜P ⎟ AR ⎝ dt dt ⎠

(19)

where AR is the anodizing ratio (=the ratio of the barrier layer thickness to the cell potential, in nm V−1), and assumed to be a constant. From eqs 16−19, the porosity (P) of porous AAO is given by P=

t +(m2 /m1) 1 − (1 − t +)(m2 /m1)

(20)

For porous AAO formed in 1.125 M oxalic acid at 30 V, Takahashi and Nagayama reported that the transport numbers of mobile Al3+ and O2− ions are t+ = 0.4 and t− = 0.6, respectively.46 3.4. Stress-Driven Ionic Transport

The high-field conduction model describes the relation between ionic current density (j) and the electric field (E) well. However, stress gradients in the oxide may possibly contribute to the ionic transport. Hebert and Houser47,48 have 7492


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to tensile transition at 0.5−1.0 mA cm−2), while average tensile stress of the order of 50 MPa was predicted above 1 mA cm−2, which is in good agreement with the experimental stress data of Bradhurst and Leach.50 In addition, by taking into consideration the viscous flow of oxide material, the model predicted the increases of the cation transport number (t+) as a function of current density (j). On the basis of experimental evidence that cation transport number (t+) is largely dependent on the electrolyte condition, Hebert and Houser pointed out that the oxide viscosity and conduction parameters may depend on the solution composition as a result of electrolyte anion incorporation into the anodic oxide film. They suggested that bulky electrolyte anions disrupt the local packing of oxygen ions and influence transport properties by the introduction of additional free volume into the amorphous oxide.48

developed a model for ionic transport in growing amorphous anodic alumina films, in which ion migration in the oxide is driven by gradients of mechanical stress as well as electric potential. It also considers the viscoelastic creep of the oxide. In other words, both stress gradient-driven ionic migration and stress gradient-driven creep are considered in the model. It is assumed that stress originates at the metal/oxide interface due to the volume change upon oxidation. For stress gradientdriven ionic migration, the empirical high-field conduction relation is generalized by considering the dependence of the ionic current density on the gradient of the ionic chemical potential ∇μi:47−49 Ji = −2

∇μi |∇μi |

⎛ a ⎞ C iu i0 sinh⎜ |∇μi |⎟ ⎝ RT ⎠

(21)

where Ji, Ci, and are, respectively, the flux, the concentration, and the pre-exponential velocity of ion “i” (i = M and O for metal and oxygen, respectively), and a is the migration jump distance in the oxide. The chemical potential μ is related to the mean stress (σ) and electrical potential (ϕ) as follows:48 u0i

μi = u i0 + z iFϕ − Vi̅ σ

4. ELECTROLYTIC BREAKDOWN When valve-metals (e.g., Al, Ta, Nb, Zr, etc.) are anodized under galvanostatic conditions, the thickness of the oxide films increases linearly with time. Correspondingly, the applied potential (U) increases linearly with time to keep the electric field (E) constant during the process. Under this condition, the anodizing potential (U) finally reaches a value at which visible sparking on the anode starts appearing, and local thickening, cracking, blistering, or even burning of oxide film commences. This local event is called “electrolytic breakdown”, which not only prevents the uniform growth of anodic films over the macroscopic metal surface, but also permanently degrades the dielectric properties of the oxide. The anodizing potential at the onset of this local event is called breakdown potential (UB). Because the oxide thickness increases linearly with the anodizing potential (U) in galvanostatic conditions, the breakdown is dependent on the oxide thickness and occurs at a critical oxide thickness. Breakdown during anodization can be associated with a number of phenomena. These include the appearance of visible sparking/luminescence,51−59 the local crystallization of oxide, 60−66 oxygen evolution at the anode,63,67,68 retardation of potential rise,69,70 occurrence of audible cracking,71 and rapid voltage fluctuations.69,70,72 In porous AAO growth, breakdown can occur under high current density anodizing conditions.4 If the reaction heat cannot be adequately dissipated from the anode, electrolyte heating may cause local increase in conductivity and a current “run away” process. This results in local thickening or burning of anodic oxide, terminating uniform growth of porous AAO. The anodic oxide in the burnt area exhibits typically a different color from the burnt-free areas. For a given anodizing electrolyte, on the other hand, porous AAO formed at a potential just below breakdown value (i.e., U < UB) exhibits the best self-ordering of pores (section 7.1).18 Improving the breakdown characteristics of anodic oxide films through proper control of the electrolyte composition, surface state of the starting aluminum, and reaction heat can allow one not only to explore new anodizing conditions for self-ordered pore growth, but also to engineer internal pore structures (see sections 7.2−7.4). In this section, we discuss some of the electrochemical factors influencing breakdown, and models that explain the breakdown phenomena.

(22)

where zi, and V̅ i are the standard chemical potential, the charge number, and the molar volume of ion i, respectively. For barrier-type anodic alumina film, the mean normal stress is defined according to σ = 1/3(σxx + σyy) = 2/3σxx, where x- and y-directions are parallel to the interface.47 For the stress gradient-driven oxide creep, the model enforces the conservation of electrical charge and volume and the momentum balance in a Newtonian fluid. For galvanostatic anodization of aluminum at the applied current density j, the constraint of charge conservation can be written as follows: u0i ,

j = −2FJO + 3FJM

(23)

On the other hand, the volume balance is jΩM = −VO̅ JO − v 3F

(24)

where ΩM is the molar volume of the Al atom in the metal and v is the creep velocity in the oxide. By employing the Maxwell viscoelastic model and also by assuming a large elastic modulus, the momentum balance in a Newtonian fluid is expressed as47 0=

1 1 ∇σ + ∇2 v + ∇(∇·v) η 3

(25)

where η is the viscosity. For porous AAO film growing under steady-state,47 the model predicted that a large compressive interfacial stress causes the lateral flow of oxide materials from the center of pore base toward the cell boundaries and the upward flow in the pore wall oxide, as in the flow pattern experimentally observed from W tracer studies (see section 6.3.5). Simulation results indicated that the stress field driving the flow results from the following three origins: “the volume expansion occurring at the metal/oxide interface, nonlinearity of the equations governing conduction of mobile ions (i.e., Al3+ and O2−/OH−), and incorporation of electrolyte-derived anionic species within the anodic oxide near the oxide/electrolyte interface”.47 For barrier-type anodic alumina film,48 the model predicted the average stress in the oxide to be compressive when the current density is smaller than 0.5 mA cm−2 (i.e., compressive

4.1. Factors Influencing Breakdown

In general, the breakdown potential (UB) is dependent on the nature of the metal being anodized, the current density (j), and the composition (or resistivity) of the electrolyte. Meanwhile, 7493


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film. For anodic oxide of aluminum, some authors have reported that anion concentration (CA−) influences UB.76,77 Kato et al. showed that at a fixed solution resistivity UB decreases linearly with an increase in the logarithm of the anion concentration, or more specifically the anion charge with the following relation:77

the electrolyte temperature, stirring rate, and history of anodic oxide have no influence on the breakdown potential. 4.1.1. The Nature of Anodized Metal. Wood and Pearson investigated metals whose anodization in 3% ammonium tartrate ended in sparking, and associated the breakdown potential (UB) with the ionic bonding characteristics of the anodic oxides by employing the criteria of Pauling and Wells. They established a descending order of UB according to the melting point of the corresponding oxide: Zr (300 V) > Al (245 V) > Ta (200 V) > Nb (190 V).72 However, Alwitt and Vijh reported a different descending order of UB for anodizations of the same metals in the same conditions: Al (350 V) > Zr (315 V) > Ta (275) > Nb (190 V).73 They correlated the increase in UB with the increasing heat of formation per equivalent (−ΔHf/equiv) of oxide, which is approximately equal to one-half the value of the forbidden band gap of the oxide. Further, they noted that the dependence of UB on the band gap would reflect the electronic nature of the breakdown phenomena. As such, rather conflicting reports have been published for the dependence of UB on the intrinsic solidstate properties of anodic oxides. Iknopisov et al.69 pointed out that the dependence of UB on the nature of the metal is considerably smaller than the dependence on the electrolyte resistivity (ρe). 4.1.2. Electrolyte Conditions. Early studies have reported that the breakdown potential (UB) increases linearly with the logarithm of the electrolyte resistivity (ρe) with the following equation:

UB = A + B log ρe

UB = A − B log C A−

(27)

On the basis of anodization experiments with tantalum in sulfuric, phosphoric, and hydrochloric acids, Arifuku et al.78 reported that UB is dependent upon the detailed distribution profiles of incorporated anions in the anodic oxide. Later, the role of incorporated electrolyte species in the electrical breakdown was emphasized by Albella et al., who have put forward a theory of avalanche breakdown during anodic oxidation.79−81 4.1.3. Current Density (j). For tantalum anodization in ammonium sulfate, Yahalom and co-workers76,82 reported that the breakdown potential (UB) is almost independent of the current density (j). For anodic films on aluminum, Ikonopisov et al.69 also reported that a 500-fold increase of current density (j) only lowers the breakdown potential (UB) by 15%. On the other hand, Di Quarto et al.74,75,83 pointed out the occurrence of two different kinds of breakdown, that is, “mechanical” and “electrical” breakdown. For anodic oxides of tungsten,74 zirconium,75 and titanium84,85 under limited conditions, they noted that anodic oxides grew with an increasing number of defects at a retarded rate (i.e., reduced slope in U−t curve) during galvanostatic anodizations, until electrical breakdown (EB) eventually occurs with visible sparks. They termed this characteristic growth as mechanical breakdown (MB). For electrical breakdown (EB), they reported that current density (j) has little effect on the breakdown potential (UEB), which is in line with the reports of Yahalom et al. and Ikonopisov et al.69,76,82 In the case of mechanical breakdown (MB), however, they observed that current density (j) has a significant effect on the breakdown potential (UMB) according to the following equation:

(26)

where A and B are the constants depending on the electrolyte composition and the anodized metal.69,71,74,75 Figure 7 shows the dependence of UB on log ρe during anodization of Nb, Ta, Al, and Zr.69 It appears from the figure that the different influences of ρe on UB defeat attempts to set the metals in series with respect to the breakdown characteristics of their anodic

UMB = AMB + BMB log j

(28)

where AMB and BMB are constants, which depend mainly on the kind of anion in the electrolyte and slightly upon pH and concentration of electrolyte: BMB > 0 for anodic oxides of zirconium and titanium75,85 and BMB < 0 for anodic oxide of tungsten.74 4.1.4. Other Factors Influencing Breakdown. The surface state of the starting metal (i.e., the surface defects (flaws), purity, processing history, etc.) also strongly influences the breakdown potential (UB).61,86 In general, the surface defects unavoidably cause a decrease of the breakdown potential (UB) with the commencement of sparks.87 On the other hand, post-breakdown anodization experiments have shown that breakdown characteristics are independent of the history of the anodic oxide film.72,88,89 When a valve-metal was anodized in electrolyte A until breakdown occurred at UB,A, and then the resulting sample was re-anodized in electrolyte B with a higher breakdown potential (UB,B), the film formation during the post-breakdown anodization continued at normal kinetics until breakdown occurred at UB,B.88,89 Temperature (T) is one of the easily controllable parameters of the electrolyte. Ikonopisov formulated the temperature dependence of the breakdown potential UB (section 4.2.1).90 However, a change in the electrolyte temperature can alter both the electrolyte

Figure 7. Dependence of the breakdown potential (UB) on the logarithm of electrolyte resistivity (ρe) for anodizations of Ta, Nb, Al, and Zr in solutions of ammonium salicyalte in dimethylformamide. Reproduced with permission from ref 69. Copyright 1979 Elsevier. 7494


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resistivity (ρe) and the property of the growing anodic oxide. When the dependence of electrolyte resistivity (ρe) is considered, no clearly pronounced dependence of the breakdown potential (UB) on the temperature (T) was obtained.71,77,91 4.2. Models for Breakdown

4.2.1. Electron Avalanche Multiplication. The first attempt to develop a quantitative model of breakdown was made by Iknopisov.90 He considered experimentally observed breakdown characteristics, and noting that the breakdown potential (UB) mainly depends on the nature of the anodized metal and the electrolyte resistivity (ρe), he inferred that breakdown is dependent upon the solid-state properties of the anodic oxide and is controlled by electrochemical reactions at the oxide/electrolyte interface. In his model, the initial electrons are injected from the electrolyte into the oxide conduction band (CB) by either a Fowler−Nordheim or a Schottky mechanism (Figure 8). The injected electrons

je, x = 0 = α1 exp[α2E1/2]

(31)

ln je, x = 0 = β1/T + β2

(32)

je, x = 0 = γ1ρe−γ2

(33)

From eqs 30 and 31, the dependence of breakdown potential (UB) on the electric field (E) is given by UB = (ψi /rq)(ln je,B − ln α1) − (ψiα2/rq) E

(34)

This equation explains a slight decrease of UB with increasing current density (j). Regarding the relation between the breakdown potential (UB) and temperature (T), the following expression is obtained by combining eqs 30 and 32: UB = (ψi /rq)(ln je,B − β1/T − β2)

(35)

Equation 35 predicts that UB is dependent on the temperature (T), which conflicts with experimental observations.71,77,91 For this discrepancy, Ikonopisov pointed out the interplay between temperature (T) and the solution resistivity (ρe). For the dependence of breakdown potential (UB) on the electrolyte resistivity (ρe), from eqs 30 and 33, one may obtain UB = (ψi /rq)(ln je,B − ln γ1 + γ2 ln ρe ) = (ψi /rq)(ln je,B − ln γ1) + (2.3ψiγ2/rq) log ρe

(36)

Equation 36 has exactly the same form as eq 26, which describes an empirical relation between UB and ρe. Although Ikonopisov’s model explains some of the experimental results, it has been criticized by many authors. Shimizu pointed out the unrealistic value of the mean free path λ(E) of ionized electrons, which from eqs 26, 29, and 36 is given by λ(E) = 1/α(E) = ψi /rqE = (1/2.3E)(B /γ2)

By using the experimental values from Ikonopisov et al. for E (=8.7 × 106 V cm−1) and B/γ2 (=1000 V), Shimizu obtained λ = 500 nm, which roughly corresponds to the thickness of oxide films formed up to the potential 400 V and indicates the absence of the electron avalanche capable of causing the breakdown.94 Albella et al. questioned the origin of electrons in Ikonopisov’s model.87 They pointed out Ikonopisov’s model lacked a reasonable explanation of the role of the electrolyte and the absence of specific electrochemical reactions required for the injection of the initial electrons. Albella et al. explicitly considered the effect of the anodizing electrolyte by posulating that the initial electrons for the avalanche come from the electrolyte species incorporated into anodic oxide.79−81 The incorporated electrolyte species act as impurity centers close to the oxide conduction band (CB), releasing electrons to the conduction band via the field-assisted Poole−Frenkel mechanism (Figure 9).80,87,95,96 In the model of Albella et al., the total current density (jt) consists of three components:

Figure 8. Schematic representation of the band structure and the avalanche breakdown in Ikonopisov’s model. Reproduced with permission from ref 79. Copyright 1984 The Electrochemical Society.

accelerate and multiply in avalanche during their travel in the oxide of thickness (tox) to the anode, until the avalanche current reaches a critical value for breakdown. In this multiplication process, the electronic current (je) depends on the travel distance (x) with x = 0 being the oxide/electrolyte interface, which can be expressed by je, x = t = je, x = 0 exp[α(E)tox ] = je, x = 0 exp[rqEtox /ψi] ox

(29)

where α(E) is the impact ionization coefficient at the electric field E, r is a recombination constant (r < 1), q is the electron charge, and ψi is the threshold energy for impact ionization. Breakdown occurs if the electronic current (je) exceeds a critical value je,B at a critical oxide thickness (tox,B). Because UB = Etox,B, the breakdown potential (UB) is given by UB = (ψi /rq)(ln je,B − ln je, x = 0 )

(37)

jt = j1 + j2 + je

(38)

where j1 is the oxidation current density, j2 is the current density consumed by the incorporated electrolyte species and is assumed to be a constant faction γ of j1 (i.e., j2 = γj1), and je is the electronic current density. The electronic current density (je) at the anode can be expressed as

(30)

The dependences of je,x=0 on the electric field (E),69,92 temperature (T), and electrolyte resistivity (ρe)93 were empirically determined to be, respectively: 7495


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V (t ) =

1 + φγ Kj t 1+γ t

ΔV (t ) =

γη E [exp(αU /E) − 1] 1+γα

(45)

(46)

The factor (1 + φγ)/(1 + γ) in eq 45 describes the correction of the anodizing rate due to the incorporation of electrolyte species. On the other hand, eq 46 enforces deviation of potential from the linearity due to the avalanche effect. Accordingly, eq 44 predicts a gradual decrease of slope (dU/ dt) of the potential−time curve during galvanostatic anodization. Albella et al. confirmed the validity of eq 44 by fitting it on the experimental results of tantalum anodization (Figure 10).80

Figure 9. Band diagram showing the avalanche multiplication of electrons in the model of Albella et al. The impurity level in conduction band (CB) is denoted by “A”. Reproduced with permission from ref 81. Copyright 1987 Elsevier.

je = je, x = 0 exp[αtox ]=je, x = 0 exp[αβU ]

(39)

where α is the impact ionization coefficient, and β is the ratio of the oxide thickness to the anodization potential (U) and is equal to the inverse of the electric field E (i.e., β = 1/E = AR, the anodizing ratio). Because the initial electronic current originates from incorporated species, je,x=0 should be a constant fraction η of oxyanion current (j2) and is je,x=0 = ηj2 = ηγj1. Under the assumption that the critical current density is a fraction z of the oxidation current j1, the breakdown potential (UB) should satisfy the following relation:

je, x = 0 exp[αβUB] = zj1

(40)

Figure 10. Experimental result for the evolution of the potential (U) as a function of time (t) during tantalum anodization in 1.2 M H3PO4 at 1.78 mA cm−2. The theoretical curve (solid line) has been fitted according to eqs 44−46. Reproduced with permission from ref 80. Copyright 1985 Elsevier.

Accordingly, the breakdown potential (UB) is given by UB = (1/αβ) ln(zj1 /j0 ) = (E /α) ln(z /ηγ)

(41)

The time derivative of the potential is given by ⎛ 1 ⎞⎛ E ⎞⎛ M M2 ⎞ dU = ⎜⎜ ⎟⎟⎜ ⎟⎜⎜ 1 j1 + j⎟ dt x 2y2 2 ⎟⎠ ⎝ ρox ⎠⎝ F ⎠⎝ x1y1

Further, by fitting eq 44 on the experimental potential evolutions in different electrolyte concentrations (C), they obtained a relation between γ and C:95

(42)

where ρox is the oxide density, F is the Faraday constant, and M1 and M2 are the molecular weights of the oxide and the incorporated species, respectively, whose corresponding anion and cation valences are x1,y1 and x2,y2, respectively. Combining eqs 38 and 42 yields the following differential equation: dU /dt = Kjt (1 + φγ )[1 + γ + γη exp(αβU )]−1

γ ≈ aC b

with a and b being electrolyte-dependent constants. From eqs 41 and 47, the concentration dependence of the breakdown potential (UB) is given by UB ≈ (E /α)[ln(z /ηa) − b ln C ]

(43)

(48)

which is in good agreement with the experiments. 4.2.2. Stress-Driven Breakdown. Sato97 distinguished five different possible contributions to the mechanical stresses in anodic oxide: (a) electrostriction pressure, (b) interfacial tension of the film, (c) internal stress caused by the volume expansion, (d) internal stress due to partial hydration/ dehydration of the anodic oxide, and (e) local stress caused by impurities. By considering the first two contributions as the most general factors for breakdown, he mathematically derived a thermodynamic model of stress-driven breakdown, and

where K is the unitary rate of anodization for oxide without electrolyte incorporation and given by M1E/x1y1ρoxF, and φ is the ratio of the equivalent weight of the incorporated species to that of oxide, that is, φ = (M2/x2y2)/(M1/x1y1). Assuming a constant field in the oxide, the integration of eq 43 yields the relation between the anodizing potential (U) and time (t): U (t ) = V (t ) − ΔV (t )

(47)

(44)

with V(t) and ΔV(t) given by 7496


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film pressure. The alignment of pores in a close-packed pattern was considered to be a process of relieving the film stress. As the breakdown in the anodic oxidation of aluminum proceeds, therefore, a porous oxide layer progressively forms and thickens on the compact barrier oxide layer, of which thickness remains constant with continuous plastic deformation. This electrostriction-stimulated breakdown model is somewhat in line with the recent flow model accounting for the steady-state formation of porous AAO film by Skeldon et al. (see section 6.3.5).98

explained the breakdown potential (UB) and the effect of adsorbed anionic species on it. According to the model,97 the stress (ΔP) accumulated in the oxide film is equal to ΔP =

ε(ε − 1) γ − 8π tox

(49)

where ε is the oxide permittivity and γ is the surface tension. In eq 49, the first term represents the electrostriction effect and the second the interfacial tension effect. According to Sato, for an oxide dielectric with ε > 10, an electric field (E) of 5 × 106 V cm−1 produces a compressive electrostriction pressure exceeding 1000 kg cm−2, which is higher than the critical mechanical strength of oxides and thus may be a cause of their mechanical failure. Because the compressive stress within the anodic oxide increases with thickness, there is a limiting oxide thickness (tox,B) above which breakdown occurs. The incorporated anionic species causes a decrease in the surface tension (γ), and thus increases the stress by lowering the value of the second term in eq 49. The dependence of the breakdown potential (UB) on the anion concentration (CA−) in the electrolyte was also established, as follows: dUB 8πkT =− ΓA− ε(ε + 1) d ln C A−

5. STRUCTURE OF POROUS ANODIC ALUMINUM OXIDE (AAO) 5.1. General Structure

Figure 12 shows schematically an idealized structure of porous AAO, together with scanning electron microscopy (SEM) images of each part of the porous AAO. Porous AAO has a honeycomb-like structure. Porous oxide layer formed on aluminum substrate contains a large number of mutually parallel pores. Each cylindrical nanopore and its surrounding oxide constitute a hexagonal cell aligned normal to the metal surface. Each nanopore at the metal/oxide interface is closed by a thin barrier oxide layer with an approximately hemispherical morphology. Under proper anodization conditions, the oxide cells are self-organized to form a hexagonally close-packed structure.7 On the other hand, the surface of the aluminum after complete removal of the porous oxide layer is textured with arrays of concave features. The thickness of the porous AAO layer on aluminum is proportional to the total charge (Qc) involved in the electrochemical oxidation. Therefore, the depth of oxide nanopores is easily tunable from a few tens of nanometers up to hundreds of micrometers by controlling anodization time (t). In general, the structure of self-ordered porous AAO is often defined by several structural parameters, such as interpore distance (Dint), pore diameter (Dp), barrier layer thickness (tb), pore wall thickness (tw), pore density (ρp), and porosity (P). For ideally ordered porous AAO, the following relationships can be drawn by simple geometric consideration:

(50)

where ΓA− is the anion density at the oxide surface at the breakdown potential (UB). The model predicts a lower breakdown potential for electrolytes having higher anion concentration. Kato et al. have also used electrostriction to explain the enhancement of breakdown by incorporated anionic impurities (see eq 27).77 They suggested that the incorporated anions lead to additional electrostrictive input into the mechanical stress in oxide films. Sato noted three different forms of mechanical breakdown depending on the mechanical property of the films: brittle crack for rigid anhydrous anodic oxides, and plastic deformation or flow for visco-plastic hydrous anodic oxides (Figure 11).97 He suggested that the formation of porous AAO films on aluminum is associated with continuous mechanical breakdown, accompanied by a continuous plastic flow of oxide under high

Dint = Dp + 2tw

(in nm)

(51)

⎛ 2 ⎞ ⎟ × 1014 cm−2 ρP = ⎜ 2 ⎝ 3 Dint ⎠

(52)

⎛ π ⎞⎛ Dp ⎞ P(%) = ⎜ ⎟⎜ ⎟ × 100 ⎝ 2 3 ⎠⎝ Dint ⎠

(53)

These structural parameters of porous AAO are known to be dependent on the anodizing conditions: the type of electrolyte, anodizing potential (U), current density (j), temperature (T), etc. Among those, anodizing potential (U) and current density (j) are the most important electrochemical parameters. A review on this matter has recently been published by Sulka.99 Here, we briefly discuss the major structural parameters of porous AAO and the electrochemical factors influencing them. 5.1.1. Pore Diameter (Dp). O’Sullivan and Wood used electron microscopy to quantitatively study the morphology of porous AAO potentiostatically formed in phosphoric acid (H3PO4) electrolyte.100 The pore diameter (Dp), interpore distance (Dint), and barrier layer thickness (tb) were observed to be directly proportional to the anodizing potential (U). Their microscopic analysis revealed that the pore diameter increases

Figure 11. Three modes of mechanical breakdown of surface films. Reproduced with permission from ref 97. Copyright 1971 Elsevier. 7497


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Figure 12. Schematic structure of (a) porous anodic aluminum oxide (AAO) on Al foil and (b) cross-sectional view. (c−e) SEM images of porous AAO, showing top surface, barrier layer, and bottom surface, respectively. Scale bars are 1 μm. Panels c−e were reprinted with permission from ref 111. Copyright 2006 Macmillan Publishers Ltd.: Nature Materials.

at the rate (ζp) of 1.29 nm V−1 with respect to the anodizing potential (U): DP = ζp·U = 1.29·U

(54)

They pointed out that the electrolyte concentration does not significantly influence the pore diameter (Dp), while the temperature of the electrolyte is positively correlated with pore diameter.100 On the other hand, theoretical modeling of porous AAO growth performed by Parkutik and Shershulsky predicted a decrease in pore diameter with decreasing electrolyte pH (i.e., increasing electrolyte concentration) due to the enhanced dissolution velocity of anodic oxide at the pore base.101 Moreover, a recent study by Sulka and Parkoła indicated that the pore diameter decreases with decreasing temperature.102 In general, pore diameter close to the surface of a porous AAO film is larger than that close to the pore bottoms (i.e., truncated pore channels), especially when anodization is conducted at an elevated temperature and/or for an extended period of time. This can be attributed to the chemical dissolution of the pore wall oxide by acid electrolyte. Accordingly, pore diameter measured from the pore bottom (not from the surface of AAO film) is more relevant for investigating the intrinsic effect of electrochemical parameters on the structure of porous AAO. Recently, Lee et al.103 reported that pore diameter (Dp) increases with current density (j) under potentiostatic anodization conditions (Figure 13). Under specific experimental conditions, they observed spontaneously oscillating current

Figure 13. Cross-section SEM micrographs of AAOs prepared from two separate anodization experiments, whose reactions were terminated near j = 86 mA cm−2 and j = 881 mA cm−2 in sinusoidally oscillating currents under potentiostatic condition (U = 200 V). (c) A schematic cross-section of AAO on Al. (d) The parameters defining the geometry of the pore bottom. Scale bars = 250 nm. Reproduced with permission from ref 103. Copyright 2010 Wiley-VCH Verlag GmbH & Co. KGaA, Weinheim.

during potentiostatic hard anodization (HA) at the potential range of 140−200 V (see section 7.2). They suggested that at a given potentiostatic condition (i.e., U = constant) the 7498


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V−1.7,100,114 For self-ordered porous AAOs formed by mild anodization (MA, see section 7.1) in oxalic and phosphoric acid at various anodizing potentials (U), Vrublevsky et al. have reported empirical equations based on results from their reanodizing experiments.115−117 They used the equations for estimating the barrier layer thickness by assuming the anodizing ratio, ARMA = 1.14 nm V−1. On the other hand, reduced anodizing ratios for sulfuric and oxalic acid have recently been reported for hard anodization (HA); ARHA = 0.6−1.0 nm V−1 (see Figure 13d and section 7.2).18,111,113,118,119 Chu et al.18 determined the anodizing ratio for less-popular anodizing electrolytes (e.g., tartaric, citric, glycolic, and malic acids) by performing what they termed “critical-potential anodization”. The anodizing ratio for various electrolytes was determined to be AR ≈ 1 nm V−1 (Figure 14). It should be noted that this value of anodizing ratio is the averaged proportionality constant determined from mild and hard anodization experiments of aluminum.

distribution of the current lines and electric field (E) may be sensitively varied by the geometric details of the barrier oxide layer. That would influence the movement rates of the electrolyte/oxide and oxide/metal interfaces, hence the pore diameter (Dp).103 5.1.2. Interpore Distance (Dint). It has long been established that the interpore distance (Dint) is also linearly proportional to the anodizing potential (U).7,100,104−109 A detailed study on this matter was performed for sulfuric and oxalic acid by Ebihara et al.104,105 Their empirical expressions on the relationship between the interpore distance (Dint) and anodizing potential (U) are as follows: Dint = 12.1 + 1.99 ·U

for sulfuric acid:

(U = 3−18 V) (55)

for oxalic acid:

Dint = 14.5 + 2.00 ·U

(U ≤ 20 V) (56)

=−1.70 + 2.81·U

(U ≥ 20 V)

(57)

For oxalic acid-based anodizations in the potential range of 20− 60 V, Hwang et al. reported that interpore distance only depends on anodizing potential (U), not on the temperature of the electrolyte:107 for oxalic acid:

Dint = −5.2 + 2.75 ·U

(U = 20−60 V) (58)

This temperature independence of the interpore distance is in line with the results of Keller et al.,7 but conflicts with the experimental results of Sulka and Parkoła,102 who observed that interpore distance is positively correlated with temperature for self-ordered porous AAOs formed by sulfuric acid-based anodization; interpore distance at an elevated temperature (10 °C) is about 10% larger than that at a low temperature (i.e., −8 to 1 °C). O’Sullivan and Wood100 reported for phosphoric acid-based anodization that increasing the temperature or the electrolyte concentration decreases the interpore distance. For self-ordered porous AAOs formed by mild anodization (MA) conditions using sulfuric, oxalic, and phosphoric acid, it has generally been accepted that the interpore distance (Dint) is linearly proportional to the anodizing potential (U) with a proportionality constant ζMA of 2.5 nm V−1 (see section 7.1):109 Dint = ζMA ·U = 2.5·U

Figure 14. Effect of anodizing potential (U) on the barrier layer thickness (tb) for porous AAO formed in different acid electrolytes. (Solid symbols, measured values; open symbols, calculated values from the half-thickness of the pore walls). Reproduced with permission from ref 18. Copyright 2006 The Electrochemical Society.

(59)

5.2. Structure of Pore Wall (Anion Incorporation)

However, this empirical formula is not valid for hard anodization (HA), under conditions in which a high electric field (E) is exerted across the barrier layer due to high current density (j) during anodization.16,110−113 This will be discussed in detail in section 7.2. 5.1.3. Barrier Layer Thickness (tb). The thickness of the barrier layer (tb) is one of the most important structural parameters of porous AAO for understanding the kinetics of the electrochemical oxidation of aluminum. Like other structural parameters, barrier layer thickness (tb) is also dependent on the anodizing potential (U). The potential dependence of the barrier layer thickness has also been known as “anodizing ratio (AR = tb/U)”, the inverse of which corresponds to the electric field (E) across the barrier layer, and it determines the ionic current density (j) (see eq 1). Accordingly, at a given anodizing potential (U), current density (j) increases exponentially as a function of the inverse of the anodizing ratio AR (i.e., the electric field strength E). Earlier studies have indicated that the anodizing ratio equals 1.2 nm

The incorporation of electrolyte-derived anions into anodic alumina is considered a general phenomenon for both barrierand porous-type anodization, occurring least for the former and greatest for the latter.3 For three major pore-forming acid electrolytes (e.g., H2SO4, H2C2O4, and H3PO4), incorporation of acid anions occurs via inward migrations under an electric field (E) during the anodization of aluminum. The incorporated acid anions influence the chemical, optical, and mechanical properties of the resulting porous AAO. For example, incorporated oxalate (C2O42−) anions together with singly ionized oxygen vacancies (F+ center) have been known to contribute to the blue photoluminescence (PL) of porous AAO formed in oxalic acid solution.120−122 The mechanical properties (e.g., hardness, wear resistance, and elasticity) of anodic alumina are also known to be affected by the incorporated chemical species (e.g., water and acid anions).123−125 The amount of incorporated acid anions and their distribution in anodic alumina depend on the anodization potential (U), current density (j), and temperature (T), as well 7499


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Figure 15. (a) Schematics illustrating the duplex structure of pore walls of porous AAO: vertical (left) and transverse (right) cross-sections. TEM plane view (b) of H3PO4-AAO and the corresponding X-ray maps of the elements: (c) phosphorus, (d) oxygen, and (e) aluminum. (f) TEM plane view of H3PO4-AAO, showing the different parts of the pore wall (i.e., the outer pore wall, cell-boundary band, and interstitial rod). Reproduced with permission from ref 135. Copyright 2009 Elsevier.

as the type and concentration of electrolytes.44,117,126−129 Accordingly, the chemical structure of the pore wall of AAOs varies with the anodization conditions. Han et al.129 have recently reported that, even at a steady-state growth condition (i.e., fixed U, j, and T), the content of anionic impurities and their incorporation depth decrease as a function of anodization time due to the progressive reduction of the electrolyte concentration, which markedly affects pore widening as well as the opening of the barrier oxide layer by wet-chemical etching. On the basis of TEM investigations of disordered porous AAOs formed in sulfuric, oxalic, phosphoric acid solutions, Thompson and co-workers suggested that pore wall oxide has a duplex structure in terms of chemical composition: an acidanion contaminated outer oxide layer next to the pores and a relatively pure inner oxide layer, as schematically shown in Figure 15a.130 TEM micrographs of porous AAO films formed in phosphoric and oxalic acids (i.e., H3PO4-AAO and H2C2O4AAO) showed a cell structure with cell-boundary bands, although the cell-boundary band in H2C2O4-AAO did not appear as well-defined as that observed in H3PO4-AAO. The cell-boundary bands became markedly more apparent upon continued exposure to the electron beam due to the preferential crystallization of the cell boundary regions. Meanwhile, the presence of a cell-boundary band was not apparent for porous AAO films formed in sulfuric acid (i.e., H2SO4-AAO). The duplex structure of the pore walls was confirmed by Ono and Masuko,131 who reported that the depth and the amount of anion incorporation in H3PO4-AAO increased linearly with anodizing potential (U), but the presence of a duplex structure was not confirmed by TEM for samples formed at U < 10 V. They also found that the crystallization rate of pore wall oxide under a strong electron-beam irradiation decreases with the increasing content of incorporated electrolyte species.132 Scanning transmission electron microscopy (STEM) and energy dispersive X-ray (EDX) point analysis on H3PO4-AAO by Thornton and Furneaux revealed that the cell-boundary

bands are composed of relatively pure alumina, whereas the material adjacent to the pores contains incorporated phosphate species from the electrolyte.133,134 Recent microscopic chemical analyses of the pore wall material of highly ordered H3PO4AAO by Le Coz et al.135 clearly indicated the presence of phosphorus-free cell-boundary bands (Figure 15b−e). The different parts of the unit cell were found to have a heterogeneous chemical composition of Al2O3·0.197AlPO4· 0.034H2O, which supports the results of the previous works by Thompson et al.10,130 The work also highlighted, as a new finding, that there is an interstitial rod material with a composition of Al2O3·0.018AlPO4·xH2O at the triple junction connecting three cells (Figure 15f).135 Thompson and Wood related the steady-state anodizing behavior of porous AAO films formed in the major anodizing acids to the distribution of the acid anions within the barrier layers and the true field strengths across the relatively pure alumina regions.14 Starting with the knowledge that the thickness ratio of the inner to outer pore wall oxide layer increases in the order sulfuric acid < oxalic acid < phosphoric acid < chromic acid, they depicted the same order of thickness ratio for the barrier oxide at the bottom of pores and correlated it to the rates of oxide formation at the same anodizing potential (U). The averaged effective electric field (E = ΔU/tb) across the barrier layer is approximately constant for the barrier layers of AAOs formed in different acid electrolytes, because the measured anodizing ratios are similar (ARMA = tb/ΔU ≈ 1.2 nm V−1). On the other hand, the potential drop (ΔU) is greater and linear across the relatively pure alumina region and smaller across the outer acid anion-contaminated region, where the potential decreases progressively toward the oxide/ electrolyte interface (Figure 16). Therefore, the true electric field across the relatively pure alumina region of the barrier layer is in the order: 7500


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H 2SO4 −AAO > H 2C2O4 −AAO > H3PO4 −AAO > H 2CrO4 −AAO

Figure 16. Distribution of the potential drop and electric field, E (slope of the voltage−distance plot), across barrier layers of porous AAOs formed in (a) sulfuric, (b) oxalic, (c) phosphoric, and (d) chromic acid. Reproduced with permission from ref 14. Copyright 1981 Macmillan Publishers Ltd.: Nature.

Figure 17. (a) The evolution of pore diameter (Dp) as a function of time (tetch) upon wet-chemical etching of porous AAOs formed in 0.3 M oxalic acid at 40 V. Wet-chemical etchings of pore wall oxide were performed in 5 wt % H3PO4 at 29 °C. The numbers in the plot are the slopes (in nm min−1) of the corresponding linear fits. (b) SEM images showing systematic increase in pore diameter (Dp) as a function of time (tetch) upon wet-chemical etching in 5 wt % H3PO4 (29 °C). Reproduced with permission from ref 129. Copyright 2013 The American Chemical Society.

Accordingly, the rate of formation of the anodic oxide films in the different acids varies with the same order given above if solid-state ionic migration across the inner layer is the ratedetermining step.14 The current density (j) is related to the potential drop (ΔU) across the barrier oxide by the high-field conduction theory (eq 1). Among the acid electrolytes above, sulfuric acid will give the highest current density (j), because of the highest potential drop across the relatively pure alumina region adjacent to the metal interface. The duplex nature of the pore wall can be experimentally evidenced by investigating the rate of pore widening. For a given set of etching conditions (i.e., temperature and concentration of an etchant solution, typically H3PO4), the rate of etching is dependent on the chemical composition of the pore wall oxide of AAO.129 Figure 17 shows (a) the evolution of pore diameter (Dp) as a function of pore wall etching time (tetch) for porous AAO formed in 0.3 M oxalic acid (H2C2O4), together with (b) representative SEM micrographs. As presented in Figure 17a, Dp versus tetch plot is characterized by an inflection point, at which the slope of the curve changes. Pore wall oxide in the early stage is etched at a higher rate (1.04 nm min−1) than that (0.36 nm min−1) in the later stage. The retarded rate of etching in the later stage can be attributed to the relatively pure nature of the inner pore wall oxide, as compared to the less dense outer pore wall oxide due to the incorporation of anionic species. As shown in Figure 17b, the

ability to precisely control the pore diameter by the pore widening process is one of the most attractive features of porous AAO for template-based nanofabrication. This feature allows one to systematically investigate the size dependence of chemical or physical properties of ordered arrays of nanodots, nanowires, or nanotube materials prepared from porous AAO templates. As-prepared porous AAOs are amorphous and contain varying amounts of water depending on anodizing condition.3 The local coordination environments of aluminum in amorphous AAOs have been extensively studied using X-ray radial distribution analysis,136 electron-yield extended X-ray absorption fine structure (EXAFS) spectroscopy,137 and magicangle spinning nuclear magnetic resonance (27Al MAS NMR).138−141 Amorphous AAOs have been considered to have a close structural relation to spinel (fcc) γ-Al2O3 with tetra- and hexa-coordinated aluminum cations in the mixing ratio of 1:2. Farnan et al.138 have reported that the coordination numbers of aluminum cations are dependent on the anodizing electrolyte: hexa-coordination for porous AAO formed in chromic acid (i.e., H2CrO4-AAOs), tetra-, penta-, and hexa7501


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thermodynamically stable hcp α-Al2O3 (corundum) is obtained above 1150 °C. Crystalline phase transformations of porous AAOs formed in different electrolytes have also been investigated by thermogravimetric (TG) and differential scanning calorimetry (DSC) analyses.139,140,144,147−151 The results have indicated that the type of anodizing electrolytes (i.e., the nature of incorporated acid anions) strongly influences transformation temperatures during heat treatment. For H2C2O4-AAOs, Mardilovich et al.144 reported that dehydration occurs up to 400 °C, dehydroxylation occurs at 400−700 °C, and incorporated oxalates pyrolyze at higher temperatures before the final transformation into thermodynamically stable α-Al2O3. A comparative study of H2SO4-, H2C2O4-, and H3PO4-AAOs by Mata-Zamora et al.150 has shown that incorporated phosphates are stable up to 1400 °C, while sulfates and oxalates decompose into gaseous SO2 and CO2 at 900 and 870 °C, respectively. Later, Kirchner et al.140 confirmed decomposition of sulfates in H2SO4-AAOs at around 900 °C by mass spectrometry (MS). In general, the initial crystallization and subsequent phase transformation of porous AAO are accompanied by changes in the morphology of pores and also in mechanical properties.139,144,151 MacQuaig et al. reported ∼6.7% loss of pore circularity, ∼15% increase in pore diameter (Dp), ∼13% decrease in pore density (ρ), and about a 2-fold increase in microhardness (from 2.5 to 4.7 GPa) upon heat-treatment of as-prepared H2C2O4-AAO up to 1200 °C.151 They attributed the changes of the overall pore structure to the densification of pore wall materials, which is associated with dehydration, dehydroxylation, and the loss of incorporated acid anions in the course of the phase transformations from amorphous to crystalline α-Al2O3. Mardilovich et al. observed sharp decreases in the flexibility of H2C2O4-AAO films at the onset of crystallization at 820−840 °C and on phase transformation from metastable ccp-alumina to hcp-alumina (i.e., α-Al2O3) at 1100−1150 °C.144 They pointed out that the fragility of porous α-Al2O3 membranes is so high that they cannot be considered for routine practical use as membranes. High temperature treatment of initially planar porous AAO membranes often leads to serious mechanical deformation (e.g., curling/rolling) or even cracking.139,144,148 Therefore, careful heat-cycling processing is required. Recently, Chang and co-workers reported that severe deformation of porous AAO membranes can be prevented by removing the acid-anion contaminated outer pore wall oxide before annealing at high temperatures.152 A proper hydrothermal treatment of H3PO4-AAOs improved crystallinity of the relatively pure inner oxide layer (i.e., the cellboundary band in Figure 15), which enhanced the etching contrast between the inner and outer oxide layers, after which the outer pore wall oxide was selectively removed from the unit cells of the AAO membranes by wet-chemical etching. Deformation-free porous α-Al2O3 membranes could be obtained by annealing of the resulting samples at 1300 °C (see Figure 19).

coordination for H2SO4- and H2C2O4-AAOs, and tetra- and penta-coordination for H3PO4-AAOs. It is now generally accepted that the aluminum exists in four-, five-, and six-fold coordination with oxygen, although the ratio of the coordination numbers varies in different samples formed in different anodizing conditions. Farnan et al.138 attributed such variations of aluminum coordination to the presence of hydroxyl groups (−OH) within alumina, with an increase of anodizing temperature favoring the hexa-coordination. Incorporated electrolyte species (e.g., acid anions, proton, and water) may be more responsible for such variations of aluminum coordination. Yet, a clear explanation for such variation of coordination number has not been given yet. 5.3. Effect of Heat Treatments

The surfaces of pore walls are hydrophilic due to the surfacebound hydroxyl (−OH) groups, and this feature allows easy modifications of the surface property via self-assembly of various functional molecules. Extensive recent reviews on this matter are given in refs 142 and 143. On the other hand, asprepared porous AAOs are highly labile to both acid and base attack. Proper high temperature heat treatments of as-prepared porous AAOs markedly improve their thermal stability and resistance against acid, base, and other corrosive chemicals, allowing the resulting AAOs to be useful as starting materials for developing various devices, which will operate in high temperatures or harsh environments. TEM and X-ray diffraction (XRD) studies have indicated that the porous AAOs undergo a series of polymorph transformations upon heat treatment up to 1500 °C in air with the following route (see Figure 18 for XRD): amorphous AAO → γ-Al2O3 → δAl2O3 → θ-Al2O3 → α-Al2O3.121,139,140,144−147 Amorphous AAOs crystallize into almost pure γ-Al2O3 at a temperature range of 820−900 °C, and then undergo successive transformations through metastable ccp δ- and θ-Al2O3 until

6. GROWTH OF POROUS ANODIC ALUMINUM OXIDE (AAO) 6.1. Stress Generation in Anodic Oxide Films

6.1.1. Volume Expansion. Oxidation of aluminum is a volume expansion process. The volume expansion during anodization can be quantitatively expressed by the Pilling− Bedworth ratio (PBR). The PBR is rigorously defined by the

Figure 18. XRD spectra of heated H3PO4-AAO (Co Kα radiation). Key: AlP = AlPO4, A = α-Al2O3, T = θ-Al2O3. Reproduced with permission from ref 139. Copyright 2005 Elsevier. 7502


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Figure 20. A schematic illustration of volume expansion during anodization of aluminum.

For three major pore-forming acid electrolytes (i.e., H2SO4, H2C2O4, and H3PO4), a wide range of volume expansion factors (kv) between 0.86 and 1.90 have been reported (Table 1).106,116,162,163,165−168 For anodizations of aluminum in H2SO4, Jessensky et al.165 investigated the relation between the volume expansion factor (kv) and the degree of selfordering of pores. They reported that optimal conditions for the best ordered arrays of pores are accompanied by a moderate volume expansion (i.e., kv = 1.22), while contraction or very strong volume expansion can result in the disordering of pores. Further, they proposed that the mechanical stress at the metal/ oxide interface, which is associated with volume expansion and increases with the anodizing potential (U), is the driving force for the formation of ordered hexagonal pore arrays, by causing repulsive forces between the neighboring pores. Li et al. reported a similar line of experimental results, noting that the volume expansion factor (kv) for the optimum anodization conditions required for ordered pore arrays should be close to 1.4, irrespective of the anodizing electrolytes (i.e., H2SO4, H2C2O4, or H3PO4).106 Later, on the other hand, Nielsch et al. proposed that the best self-ordering of pores requires a porosity (P) of 10%, the condition that corresponds to a volume expansion of porous AAO of about 1.23, independent of the electrolyte.109 More systematic investigations on the effect of anodization conditions on the volume expansion factor (kv) have been performed by Vrublevsky and co-workers. For porous AAO growths in H2SO4 and H2C2O4 electrolytes, they determined that the volume expansion factor (kv) has a linear dependence on the anodizing potential (U) and also noted a linear relation between the logarithm of the current density (i.e., ln j) and the inverse volume expansion factor (i.e., 1/kv).116,163,167 For anodizations under galvanostatic conditions, an increase of temperature led to a decrease of the volume expansion factor (kv) due to the corresponding decrease of the formation potential (U), in which the slopes of ln j versus 1/kv curves are invariant with respect to the anodizing temperature.163,167 However, the slopes of curves were found to be different for different electrolytes, which was explained by the influence of acid-anion incorporation on the volume expansion: the larger is the amount of incorporated acid anions, the larger is the volume expansion factor (kv).167 On the basis of the experimental results of Vrublevsky et al., one may obtain the following relation:

Figure 19. A photograph of the porous α-Al2O3 membranes obtained by annealing the porous AAOs with (a) and without (b) acid-anion contaminated outer pore walls. (c−e) SEM images of porous Al2O3 membranes obtained by annealing the porous AAOs after removal of outer pore walls at (c) 1115 °C, (d) 1250 °C, and (e) 1300 °C; the scale bars are 1 μm. Reproduced with permission from ref 152. Copyright 2012 The Royal Society of Chemistry.

molar volume ratio of grown oxide (Vox) to the consumed metal (Vm) as follows: PBR =

Mox ρm Vox = Vm nM mρox

(60)

where Mox is the molecular weight of oxide, Mm is the atomic weight of metal, n is the number of atoms of metal per one formula of the oxide, and ρm and ρox are the densities of metal and oxide, respectively. In science on the corrosion of metals, PBR has been the basis for judging the protectiveness of a passivating oxide: if PBR < 1, the passivating oxide is under tensile stress and easily cracked; if 1 < PBR < 2, the oxide covers the metal uniformly and is protective; if PBR > 2, the passivating oxide is under too much compressive stress and easily crumbles (e.g., iron oxide on iron).153,154 For anodic alumina growth, PBR can be experimentally determined from the current efficiency (ηj) of oxide formation and the density (ρAAO) of the resulting AAO, provided that the composition of anodic oxide is well-defined. The densities of barrier- and porous-type AAOs have been reported to be in the range of ρAAO = 2.7−3.5 g cm−3.104,111,155−157 Assuming composition stoichiometry of Al2O3 and ρAAO = 3.0 g cm−3, PBR for AAO growth is 1.70 at 100% current efficiency (ηj). For porous-type AAO growth, on the other hand, PBR can vary between 1.02 and 1.58 due to the lower current efficiency (ηj = 60− 93%).158−161 The exact determination of PBR is rather complicated by the composition of anodic oxide, its density, and the current efficiency (ηj) of oxide formation. Consequently, volume expansion in anodic oxide growth has been considered instead, by employing a volume expansion factor (kv), and a simple thicknesses ratio:116,162−164

k v = hAAO/hAl

6.1.2. Stress Measurements. Anodic alumina is a dielectric material. During anodization, a very large electric field (typically, 106−107 V cm−1) is impressed on the oxide film. The electric field drives the inward movement of O2− ions and the outward migration of Al3+ ions within the anodic oxide. The resistance to these counter-migrations and the attraction of

(61)

where hAAO and hAl are the vertical heights of the AAO and the consumed aluminum, respectively (Figure 20). 7503


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Table 1. Volume Expansion Factors (kv) Reported for Porous Anodic Aluminum Oxide (AAO) Growth in Three Major PoreForming Acid Electrolytes electrolyte

concentration

H2SO4

20 wt % 20 wt % 1.7 wt % 1.7 wt % 10 wt % mixture solutiond

temp (°C) 1 1 1 10 18−22 0

H2C2O4

2.7 wt % 0.34 M 0.22−0.92 M 0.1−2.9 M

H3PO4

1 15 16−18 18

Ua or jb

kvc

refs

18−25 V 19 V 25 V 25 V 6−20 mA cm−2 50 V 100 V 300 V 40 V 40 V 1.62−4.8 mA cm−2 2−13 mA cm−2

0.86−1.62 1.41 1.36 1.40 1.4−1.6 1.72 1.76 1.90 1.42 1.18 1.25−1.42 1.2−1.62

165 106 106 106 167 162 162 162 106 168 163,167 166

U = anodizing potential (in V). bj = current density (in mA cm−2). ckv = volume expansion factor. dA mixture solution of 0.41 M H2SO4 and 0.16 M H3BO3.

a

and Leach50 conducted deflection measurements on aluminum and reported the effects of current density and oxide thickness on the stress. During anodization of aluminum in ammonium borate and ammonium citrate solutions, they observed that the stress can be either tensile or compressive depending on the current density. Compressive stress was observed below 1 mA cm−2, while tensile stress was observed at higher current densities. A similar line of results was obtained by Nelson and Oriani,169 who performed deflection measurements during anodization of aluminum and titanium in 0.1 M H2SO4. The deflection was related to the stress as follows:169

the charged species result in a compressive electrostatic stress along the direction of the electric field. The compressive stress normal to the oxide surface is proportional to the dielectric constant (ε) and the square of the electric field (E):169 ε 2 σ⊥ = E (62) 8π This compressive stress occurs only under the electric field. Thus, the stress will relax when the anodization is stopped. Proost et al. have determined electrostatic stress in anodic oxide films of thickness tb as follows:67,170 σ ES = −

ε − (α1 + α2) E2 v ε0 1−v 2 tb

σ =

(63)

where v is the Poisson coefficient of the oxide film, ε is its relative dielectric constant, εo is the vacuum permittivity, and α1 and α2 are both characteristic electrostriction constants. Porous AAO has been known to incorporate water as well as varying amounts of electrolyte-derived anionic species (see section 5.2). Compositional effects arising from anion incorporation and oxide inhomogeneities may be sources of stress in anodic oxides. The stress during anodization can be measured by monitoring the changes of the substrate deflection. Stoney related the measured deflection (i.e., the radius of curvature) to the stress (σ) as follows:171

σ=

Yt 2 6Δktb

EMt 2Δk 2 3(1 − vM )L2tb

(65)

where EM and vM are the elastic modulus and Poisson’s ratio of metal strip of length, L, and thickness, t; tb is the thickness of the oxide layer where the stress is generated; Δk is the radius of curvature of the metal strip. The compressive stress due to electrostriction was found to increase linearly with anodizing potential during oxide growth. Slowly grown oxides have greater electrostrictive deflections than more rapidly grown oxides. For aluminum anodized in 0.1 M sulfuric acid, the deflections are compressive at low current densities and become tensile above 0.6 mA cm−2. The development of tensile stress was attributed to the volume difference between the metal being oxidized and the oxide formed at the metal/ oxide interface.169 Moon and Pyun investigated the effect of electrolyte concentration and current density on the deflection behavior of aluminum in sulfuric acid.173,174 Their experiments found that the deflections become more tensile as the current density increases. Compressive stresses were always observed at the relatively low current density (j = 2 mA cm−2). On the other hand, at larger concentrations of sulfuric acid, the rates of deflection with respect to the current density became nonlinear. The observed compressive stress at low current densities and tensile stress at higher current densities was explained in terms of the annihilation of cation vacancies and the formation of oxygen vacancies at the metal/oxide interface, respectively. The authors suggested that the stresses that developed are not distributed over the entire oxide film, but are limited to a narrow region of the metal/oxide interface below 1 nm.174 An opposite evolution of internal stress has recently been reported. Proost et al.175 measured the internal stress of porous AAO

(64)

where Y is Young’s modulus of the substrate metal, t is the thickness of the metal foil, tb is the thickness of the oxide on the metal, and Δk is the radius of curvature. Equation 64 has been modified over the decades to account for lateral strain, differences in the elastic moduli of the oxide and its metal substrate, and non-uniformity of the stress distribution in the oxide. Vermilya172 applied the Stoney method for anodizing different metals. Upon applying potential, the substrate deflected, indicating tensile stress. The stress observed with the forming voltage applied was always more compressive than at zero voltage except for tungsten. Higher stress was observed when the rate of oxide formation was high. Vermilya attributed the observed stress to a dynamic hydration process. As the oxide film is buried by newly generated oxide, it is dehydrated by proton migration, producing tension in the film. Bradhurst 7504


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during the very initial stage of pore formation in 0.4 M phosphoric acid by employing a high-resolution in situ curvature measurement technique based on multiple-beam deflectometry. The internal stress was found to remain constant during the initial barrier growth. However, the internal stress value increases in the compressive direction with increasing current density. When the current density is lower than 4 mA cm−2, the internal stress in the barrier oxide is tensile, while at higher current densities, the internal stress becomes compressive. The different stress evolutions in different electrolyte systems have not yet been fully understood. More systematic in situ measurements of stress as a function of the anodizing conditions at the very initial stage of pore formation are required to resolve conflicting stress evolutions reported in the literature. A meaningful advance in this direction has recently been made by Hebert and co-workers,176 who demonstrated the effectiveness of phase-shifting curvature interferometry as a new technique for high-resolution in situ monitoring of stress evolution during anodization. The newly developed metrology allows extremely stable and reliable measurement of curvature changes at a curvature resolution of 10−3 km−1, which is comparable to or higher than that of the high-resolution multiple-beam deflectometry technique.175,177−179 From stress measurements during galvanostatic anodizing of aluminum in 0.4 M phosphoric acid, Hebert and co-workers176 found that the apparent stress in the barrier oxide is tensile (+100 MPa) at low current density but became increasingly compressive at higher current densities. In addition, they observed that transition from tensile to compressive stress occurs at current density of 4 mA cm−2, in good agreement with the report of Proost and co-workers.175 6.1.3. Effects of External Stresses on Pore Growth. The effect of an applied external tensile stress on the self-ordering of pores was first studied by Sulka and co-workers.180 It was found that the magnitude of applied tensile stress influences the ordering degree of pores. At a relatively low external tensile stress, regular hexagonal arrangement of pores was observed. However, a high tensile stress completely destroyed the pore arrangement of porous AAO. Large holes and pits, rather than nanopores, appeared on a highly stressed surface. This experimental result revealed that the mutually repulsive mechanical force between neighboring pores, which is associated with volume expansion due to oxidation of the metal, may be the driving force for the self-organized formation of hexagonally close-packed arrangement of nanopores, as suggested by Jessensky et al.165 As to the effects of compressive stress on pore growth, Park et al.181 have performed the anodization of aluminum confined in micrometer-sized vertical trench patterns. The authors observed that during the anodization of the confined aluminum, the anodization rate is significantly retarded at the vertical sidewall of the trench. Because of the retarded anodization rate, most of the aluminum at the edge part of the structure remains and its thickness decreases gradually with increasing distance from the vertical sidewall toward the central part of the trench structure (Figure 21). The authors attributed this phenomenon to the accumulation of compressive stress at the vertical sidewall of the trench structure, where linear vertical volume expansion is severely prohibited by additional stress. The authors noted that because compressive stress is an additional kinetic barrier to the electrochemical oxidation of aluminum, the anodization kinetics of aluminum should be severely retarded.

Figure 21. Cross-sectional SEM images of the AAO confined in a micrometer-sized vertical trench pattern: (a) low and (b) high magnification image. Scale bar = 1 μm. Reproduced with permission from ref 181. Copyright 2006 The Electrochemical Society.

Retardation of anodization rate has also been reported for lateral anodization processes. In that process, an aluminum thin film deposited on a substrate is sandwiched by a rigid insulating top layer, and then the side edges of the aluminum are anodized to produce horizontal arrays of pores on the substrate.182−186 Oh and Thompson186 reported abnormal behavior in the anodic oxidation of aluminum in mechanically confined structures used for the formation of horizontal nanoporous AAO. The authors observed that dendrites, periodic internal pore structures (see Figure 22a), formed with a 5% retarded growth rate, as compared to its value during bulk anodization under the same conditions. They attributed the observed anomalies to the suppressed volume expansion and a plastic flow of anodic oxide confined by an insulating top layer; because volume expansion by plastic flow in the pore growth

Figure 22. (a) Cross-sectional SEM image of AAO nanopore formed by horizontal anodization. Scale bar = 100 nm. (b) Schematic illustration of the formation mechanism of horizontal AAO with dendritic internal pore structure. Reproduced with permission from ref 186. Copyright 2011 The Electrochemical Society. 7505


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Figure 23. Schematic diagram of the kinetics of porous AAO growth in (a) potentiostatic and (b) galvanostatic conditions: (a) Current (j)−time (t) curves for potentiostatic anodization (i.e., U = constant) and (b) potential (U)−time (t) curve for galvanostatic anodization (i.e., j = constant).

maximum due to the ready diffusion of electrolyte (stage III). After that, current (j) reaches a steady value after passing through an overshoot (stage IV). The appearance of current overshoot has been related to the decrease of the initial pore density with the steady-state growth of major pores:168 pores increase in size by persistent merging with adjacent pores. For a given set of anodization conditions, the rate of potential increase at the very beginning of anodization, the value of the minimum current, the time needed for anodizing current to reach a steady value, and the appearance of the current overshoot have been known to be directly dependent on the anodizing potential (U), electrolyte pH and temperature, and the initial surface state of the aluminum.168,187,188 For the case of galvanostatic anodization, a similar progression can be observed for stages I−IV, while the potential (U) changes as a function of time (Figure 23b). Under constant-current conditions, the oxide growth rate should be proportional to the applied current density (j) and constant according to the Faraday’s law. In addition, a constant electric field (E = U/tb) is required to sustain the applied constant current (j).76 Accordingly, the potential (U) increases with the thickness of the growing barrier oxide (tb), as shown in the inset of Figure 23b. However, in practice, the evolution of potential (U) deviates from a simple linear increase with time, as shown in Figure 23b. For convenience, various mechanisms governing such a deviation have been referred to as growth instabilities, which include, for example, mechanical breakdown during zirconium anodization, and surface undulation/pore initiation during aluminum anodization.178 Figure 23b shows a gradual retardation of potential (U) increase at stage II. Such a potential evolution can be attributed to a morphological instability, that is, transition from the stage of barrier oxide growth to the stage of porous oxide growth.178 In the following sections, we will discuss in detail the kinetics and morphological instability involved in the early stage of anodization, which have been systematically investigated by Proost and co-workers.160,175,178,189,190 6.2.2. Kinetics of Porosity Initiation. Recent publications have shown that the growth of a porous AAO and its selforganization are most likely driven by the internal stresses developing in the anodic oxide.47,190−192 Proost and co-workers have investigated the initiation of porosity during galvanostatic

direction is prohibited by traction at the insulating top layer, the extra volume of newly formed anodic alumina is extruded inside the primary pores, resulting in a dendritic structure, as schematically illustrated in Figure 22b. 6.2. Initial-Stage Pore Formation

6.2.1. Qualitative Description on Pore Formation. Porous AAO can be easily fabricated by anodization of aluminum in acid electrolytes either under a constant potential (i.e., potentiostatic) or a constant current (i.e., galvanostatic) condition. In general, potentiostatic anodization is widely employed for the fabrication of self-ordered porous AAO, because of the linear relation between the applied potential (U) and the structural parameters of the resulting AAO (i.e., pore diameter Dp, interpore distance Dint, and barrier layer thickness tb, section 5). Figure 23 shows (a) a typical current (j)−time (t) curve for potentiostatic anodization, (b) potential (U)−time (t) curve for galvanostatic anodization, together with (c) schematic illustrations of the stages of porous structure development.101 When a constant anodic potential (U) is applied, a thin compact barrier oxide starts to grow over the entire aluminum surface (stage I). Thickening of the initial barrier oxide over time (t) results in an increase of the series resistance (R) of the anodization circuit. Current (j) is initially maintained at the limiting current (jlimit) of the power supply, and correspondingly potential (U = jR) increases linearly with time (t) (see the inset of Figure 23a). When the thickness (or the resistance, R) of the compact barrier oxide layer reaches a certain value, current (j) drops rapidly to hit the minimum value (stage II). For this stage, O’Sullivan and Wood100 suggested that current (i.e., electric field) concentrates on local imperfections (e.g., defects, impurity, pits) existing on the initial barrier oxide, resulting in non-uniform oxide thickening and pore initiation at the thinner oxide areas. Thompson and co-workers10,33,44 have proposed that local cracking of the initial barrier oxide due to accumulated tensile stress (PBR < 1) may develop the paths for electrolyte penetration. Local increase in field strength at the penetration paths effectively polarizes the Al−O bonds, facilitating field-assisted oxide dissolution there (section 6.3.2),100 and eventually leads to development of individual penetration paths into embryo pores.10,44 Accordingly, further anodization leads to a gradual increase in current (j) to a local 7506


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Figure 24. (a) Curvature change and cell voltage evolution as a function of anodization time at 5 mA cm−2 in 1 M H2SO4 (22 °C); (b) radial power spectral density (PSD) of samples anodized for different times. Reprinted with permission from ref 190. Copyright 2009 AIP Publishing LLC.

anodization of aluminum thin film on a silicon substrate in 1 M sulfuric acid (22 °C) at j = 5.0 mA cm−2 by in situ monitoring of the internal stress-induced curvature of the substrate.190 They observed that the magnitude of the substrate curvature (ΔK, in km−1) increases at a constant rate in the compressive direction up to a transition time (ca. 7.9 s) at which the rate of curvature change suddenly increases (Figure 24a). The observed curvature transition was attributed to the initiation of porosity accompanied by an increase of the oxide growth rate. This was confirmed by a quantitative analysis of the radial power spectral density (PSD) distributions of the anodized surfaces as presented in Figure 24b, in which the first (λ1) and second (λ2) correlation breaks are associated with the aluminum grain morphology, and the appearance of porosity, respectively. Further, the authors pointed out that the increase in growth rate after the curvature transition would not be expected if the porous layer grew by dissolution of the oxide at the pore base.190 Hoar and Yahalom76,187 have proposed that the pore initiation occurs at a sufficiently low electric field due to proton entry into the initial barrier oxide at preferred sites, where concentrated electric field (i.e., current) accelerates oxide dissolution to develop pores. The authors187,193 suggested that the Al3+ ions found in the solution originated from the oxide at the pore base as a result of field-assisted (rather than “thermal”) oxide dissolution (see section 6.3.2). Yet later studies29,30,161 have revealed that Al3+ ions in the electrolyte are the result of direct ejection of Al3+ ions from the metal/oxide interface through the oxide into the solution (see section 6.3.4), rather than coming from the field-assisted oxide dissolution process. Recently, Proost et al. have for the first time correlated the kinetics of Al3+ loss to the morphological changes occurring during the very early stage of galvanostatic anodization of aluminum in 1 M sulfuric acid160 and 0.4 M phosphoric acid189 by employing in situ inductively coupled plasma optical emission spectrometry (ICP-OES). For both anodization cases, the authors observed three distinct regimes of Al3+ loss. However, the evolution of the Al3+ loss rate turned out to be markedly different, as follows. For the same current density, the rate of Al3+ loss is higher in phosphoric acid than in sulfuric acid during the barrier layer growth stage (regime A, stage I in Figure 23b). On the other hand, the Al3+ loss rate is lower during the pore initiation stage (regime B, stage II in Figure 23b), as compared to the barrier oxide layer formation stage

(i.e., regime A). This implies that the oxide formation efficiency (ηj) during the initiation of porosity is higher than that during the initial barrier oxide formation, which is in line with the increase in the oxide growth rate upon pore initiation, as shown by in situ monitoring of the internal stress-induced curvature of the substrate (vide supra).190 On the other hand, the authors attributed the lower Al3+ loss rate during the pore initiation stage to the non-uniformity of current distribution upon commencement of pore initiation.189 In sulfuric acid, the Al3+ loss rate during stages II and III was constant, and increased slightly at the beginning of stage IV. In phosphoric acid, however, the cation loss rate decreased again during stage II and then remained more or less constant during stages III and IV (regime C). The authors associated the difference in the observed evolution of Al3+ loss rate with the morphological differences of the growing anodic oxide (i.e., pore size and spacing).189 In sulfuric acid, the rate of Al3+ loss during the steady-state pore growth stage was similar to the level of the barrier layer growth stage, whereas in phosphoric acid both rates were markedly different; there was a higher rate during the barrier layer growth stage. For both cases of anodization, the rates of Al3+ loss were observed to linearly increase with the current density (j). Proost et al.189 suggested that the direct proportionality between the rate of Al3+ loss and the current density (j) can be attributed to the direct ejection of Al3+ ions into the electrolyte, because the field-assisted contribution may be relatively independent of the current density according to the high field conduction theory.25 For both cases of anodizations, the oxide formation efficiency (ηj) during the porous oxide growth stage was found to increase with current density (j). For phosphoric acid anodizing, however, the efficiency (ηj) during the barrier oxide growth stage did not increase with current density (j), and remained constant in the current range j = 2.0−10 mA cm−2 (Figure 25), which is different from the case of sulfuric acid anodizing.160 The abnormal efficiency values at 1.0 mA cm−2 in Figure 25 were suggested to occur when field-assisted dissolution (see section 6.3.2) rather than direct cation ejection (see section 6.3.4) becomes the predominant mechanism of Al3+ loss for current densities lower than 2.0 mA cm−2.164,189 6.2.3. Morphological Instability. As was briefly mentioned at the end of section 6.2.1, at the very early stage of anodization, pores initiate as a result of the morphological instability of growing anodic oxide. This becomes evident at an 7507


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(i.e., potential vs time) was characterized by three well-defined regimes, as discussed in the previous section. By converting the curvature versus time curve (see Figure 24a) into a stress− thickness versus potential curve, the authors obtained internal stress data corresponding to the barrier layer growth regime from the slope of the stress−thickness versus potential curve (Figure 27). On the basis of the anodizing ratio (AR) of 1.1 nm

Figure 25. Dependence of the oxide formation efficiency (ηj) on current density (j) for Al anodizing in 0.4 M H3PO4 during barrier growth (regime A, i.e., stage I in Figure 23b, ■) and steady-state porous growth (regime C, i.e., stages III−IV in Figure 23b, □). Reproduced with permission from ref 189. Copyright 2011 The Electrochemical Society. Figure 27. (a) Stress−thickness versus cell potential curves during galvanostatic anodization of aluminum for 20 s in 1.0 M H2SO4 (22 °C) at 5 mA cm−2. The vertical dashed line indicates the moment of pore initiation. (b) Oxide thickness at pore initiation versus instantaneous internal stress in the barrier. The numbers near data points indicate corresponding current density in mA cm−2. Reprinted with permission from ref 178. Copyright 2011 Elsevier.

oxide thickness of a few nanometers as shown in Figure 26.67,161,168,178,190,194 Previous research has suggested that

V−1 for aluminum anodizing in 1.0 M H2SO4, they then could plot the thickness at pore initiation (the first vertical dashed line in Figure 27a) as a function of stress for a set of anodizing experiments performed at different current densities, j = 1.5−25 mA cm−2 (Figure 27b). It turned out that oxide thickness at the point of pore initiation increases with the compressive stress. The authors claimed that such a dependency revealed that “pore initiation itself does not correspond to a stress-affected instability, although further development of the instability into an ordered steady-state pore morphology can be considered to be stress-affected (i.e., porosity-induced stress relaxation).”178 Instead, as an alternative instability criterion, they suggested that electrostatic energy acts as a driving force not only for porosity initiation, but also for selection of the interpore distance (Dint) of anodic oxides.175 More recently, Hebert and co-workers194 performed a linear stability analysis of an instability mechanism controlled by oxide dissolution and ionic migration at the initial stage of pore formation. The authors claimed that previous models101,192,196−198 based on nonlinear interface kinetics may be unrealistic, because the oxide formation efficiency (ηj) is weakly dependent on the current density (j).199,200 Their model predicted that the range of oxide formation efficiencies (ηj) producing pattern selection depends on the cation charge (z) and PBR; patterns with a minimum pore spacing occur within a narrow range of the oxide formation efficiency (ηj = 65−70% for porous anodic alumina and 50−58% for anodic titania), which occurs if z > 2. According to the model, the wavelength for the maximum disturbance growth rate is proportional to the thickness of anodic oxide, which quantitatively explains the proportionality of interpore distance (Dint) to anodizing potential (U). This holds for both disordered and self-ordered porous AAOs, and also for diverse anodizing electrolytes.

Figure 26. Cross-sectional TEM micrographs of anodic oxide film on Al anodized in 0.4 M H3PO4 at 4.5 mA cm−2 for (a) 17 s, (b) 34 s, and (c) 55 s. Adapted with permission from ref 161. Copyright 2010 The Electrochemical Society.

internal stress may play a role in the growth instability of anodic oxides.64,65,75,192,195 Van Overmeer and Proost178 have investigated the relation between the internal stress, the morphological instability, and the pore initiation during the growth of porous AAO. They employed a high-resolution in situ curvature measurement technique to monitor the internal stress during anodization of a thin aluminum film on a silicon substrate in 1.0 M H2SO4 at 5 mA cm−2. The anodization curve 7508


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anodization. Hoar et al.187,193 suggested that the local oxide dissolution is assisted by the increased electric field (E) due to the geometry of the pore base. The proposed model suffers disadvantages, however, in that it does not specify the details associated with field-assisted oxide dissolution and it adopts the idea that the outwardly migrating Al3+ ions contribute to the oxide formation at the oxide/electrolyte interface, which actually does not occur in the case of porous-type oxide growth (see section 6.3.4).29,30 O’Sullivan and Wood100 proposed a detailed physical mechanism for field-assisted chemical dissolution of anodic oxide, and qualitatively explained the dependence of the porous morphology on anodizing conditions. They explained the fieldassisted oxide dissolution in terms of the effective polarization of Al−O bonds in the lattice at the oxide electrolyte interface under the field (Figure 28).100 In this model, the electric field

6.3. Steady-State Pore Formation

6.3.1. Joule’s Heat-Induced Chemical Dissolution. For steady-state pore growth, the movement rates of the metal/ oxide interface and the oxide/electrolyte interface should be balanced, keeping the barrier layer thickness (tb) constant. It has long been believed that this balance is achieved by a dynamic equilibrium between the oxide formation at the metal/ oxide interface and the removal of oxide at the oxide/ electrolyte interface.100,101,168,193 Joule’s heat-induced chemical dissolution of the barrier oxide by acid electrolyte has been suggested.7,127,168,187,201 However, this thermal mechanism does not reasonably explain the dynamic balance of the movement rates of the interfaces (i.e., the metal/oxide and oxide/electrolyte interfaces), because the rate of chemical dissolution is typically much lower than that of oxide formation even at an elevated temperature.29 Generated heat has been considered to play a role in enlarging pores by assisting dissolution of the pore wall oxide, resulting in truncated pore channels.202 From the experimentally determined oxide formation rate (372.5 nm/min) in 1.5 M H2SO4 (21 °C) at 20 mA cm−2 and the chemical dissolution rate (0.084 nm/min) in the same electrolyte, Hunter and Fowle201 inferred that the electrolyte condition at the pore base must locally change to the equivalent of boiling of 5.10 M H2SO4 at 124 °C to satisfy the required rate of chemical dissolution (i.e., 372.5 nm/min). On the other hand, on the basis of their calculations of steady-state temperature distribution, Nagayama and Tamura203 reported that local temperature rise due to Joule heating under a comparable anodization condition (i.e., 1.0 M H2SO4, 27 °C, and j = 9.4 mA cm−2) is not more than ΔT = 0.07 °C. They claimed that the high rate of local oxide dissolution should be interpreted as a consequence of high electric field (E) impressed on the barrier layer, as had been suggested by Hoar and Mott.193 Recent in situ measurements of anode temperature during anodizations in 0.3 M H2C2O4 at 40 V have shown that the maximum temperature change is ΔT ≈ 1 °C, which contradicts Joule’s heat-induced chemical dissolution of anodic oxide at the pore base.204 6.3.2. Field-Assisted Oxide Dissolution. Hoar and Mott proposed193 that the thickness of the barrier oxide layer is maintained by the dynamic rate balance between the following two processes occurring at the oxide electrolyte interface: (1) the oxide formation by the reaction between O2− ions and Al3+ ions migrated from the metal/oxide interface, as in the formation of barrier-type oxide, and (2) the oxide dissolution. The authors assumed that oxide formation takes place both at the metal/oxide interface and at the oxide/electrolyte interface, as in the barrier-type oxide formation. In their model, oxide at the oxide/electrolyte interface is decomposed to Al3+ and O2− ions. The resulting Al3+ ions go into the electrolyte. Meanwhile, O2− ions in contact with acid electrolyte become OH− and move through the oxide to form new oxide at the metal/oxide interface. The proton (H+) released by the oxide formation reaction would then diffuse back to the electrolyte by proton transfer between the lattice O2− ions. Since the oxide ions from the oxide/electrolyte interface are spread over a larger area at the metal/oxide interface, the oxide dissolution occurs at a greater rate compared to oxide formation at the metal/oxide interface. In other words, the net result of process (2) is the progressive thinning of the barrier oxide layer due to the requirement of oxygen volume conservation in the oxide, which is compensated by the oxide formation through process (1) to keep the thickness of the barrier layer constant during

Figure 28. Schematics of field-assisted dissolution mechanism by O’Sullivan and Wood: (a) before polarization, (b) after polarization, (c) removal of Al3+ and O2− ions, and (d) remaining oxide. Adapted with permission from ref 100. Copyright 1970 Royal Society Publishing.

(E) across the barrier oxide can effectively polarize (i.e., stretch) the Al−O bond along the applied field direction, lowering the effective activation energy for bond dissociation. Solvation of Al3+ ions by water molecules via activated complex (i.e., Al(H2O)63+) and the removal of O2− ions by H3O+ ions as H 2O are facilitated. Because the electric field (E) is concentrated on the pore base, the oxide dissolution rate is the greatest there, and a dynamic equilibrium between the oxide formation and the oxide dissolution can be established. This model has popularly been cited in the literature to explain the growth and morphology of porous AAO. Several models based on field-assisted oxide dissolution have been developed. In their theoretical modeling, Parkhutik and Shershulsky101 considered the three-dimensional (3D) distribution of electric field and current in the barrier oxide layer and included the field-assisted dissolution at the oxide/electrolyte and oxide formation at the metal/oxide interface as boundary conditions to predict the steady-state pore morphology. Their model predicted that the movement rate of the oxide/ electrolyte interface by field-assisted dissolution is pH-dependent and has an exponential dependence on the electric field (E). 7509


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The model also predicted a linear dependence of interpore distance (Dint) on anodizing potential (U). Thamida and Chang198 were also able to predict the linear dependence of the interpore distance (Dint) on the anodizing potential (U). They performed a perturbation analysis of the equation describing the movements of the metal/oxide interface by oxide formation, and oxide/electrolyte interface by oxide dissolution, the rate of which was considered to be governed by the local electric field. The electric field at both interfaces was considered to be dependent on the shape or topography of the interfaces. Their model predicted that the ratio of pore diameter (Dp) to interpore distance (Dint) is a factor independent of anodizing potential (U), but varies with the electrolyte pH.198 Singh et al.192 have further put forward previous models. The main element of their model is the Butler−Volmer relation, which describes the exponential dependence of the current on the overpotential and the dependence of the activation energy of oxide dissolution reaction on the Laplace pressure and the elastic stress within anodic oxide.192,205 Their model predicted ordered hexagonal pore arrays for a range of volume expansion factors close to the one observed for ordered porous films. Recently, Friedman et al.206 performed a systematic experimental investigation to study the stability phase diagram as a function of pH and anodizing potential (U) in an attempt to validate the above-mentioned theoretical models. By considering the discrepancies between the previous theoretical models and their experimental results, they concluded that the previous models must include an appropriate weighting factor to account for the oxide formation and dissolution mechanism during the pore formation. The discrepancies may originate from the fact that these models were based on pore formation by field-assisted oxide dissolution at the pore base. In fact, the field-assisted dissolution mechanism for steady-state pore formation has recently been rejected by several authors (see sections 6.3.4 and 6.3.5). For the initial stage of anodization, on the other hand, Oh and Thompson207 have recently reported direct experimental evidence of the impact of the electric field (E) on the oxide dissolution rate and the existence of a threshold electric field (E*) for pore initiation. To assess the effect of an electric field (E) on the Al2O3 dissolution rate, the authors re-anodized a planar pre-formed Al2O3 layer (thickness h0 = 160 nm) on aluminum and investigated the evolution of thickness of the preformed oxide layer as a function of time (Figure 29). The thickness of the planar barrier oxide was found to have decreased from 160 to 131 nm after re-anodization in 5 wt % H3PO4 at 86 V for 49 min, without changing the thickness of the aluminum (i.e., no anodic oxidation of aluminum). In addition, the dissolution rate of the barrier oxide of a fixed initial thickness was found to increase with electric field (E). The results indicated that field-assisted oxide dissolution is operative at the initial stage of anodization. On the basis of the invariance of the thickness of aluminum and the morphology of the metal/oxide interface, Oh and Thompson suggested that the formation of the incipient pores is associated with fieldinduced instability at the oxide/electrolyte interface at sufficiently high electric field. In a separate work, Skeldon and co-workers have reported confirming experimental results of field-assisted oxide dissolution at the initial stage of pore formation during potentiostatic anodization of aluminum in phosphoric acid.208 In their study of pore initiation, the authors employed immobile arsenic species as tracers. 18O-labeled

Figure 29. (a,b) SEM images showing field-assisted dissolution of anodic oxide under an electric field (E): (a) before and (b) after reanodization of a planar preformed Al2O3 layer in a 5 wt % H3PO4 solution at 86 V for 49 min. Scale bar = 200 nm. (c) Changes in the thickness of a planar preformed Al2O3 layer due to electric field- and time-dependent dissolution behavior in a 5 wt % H3PO4 solution. Reprinted with permission from ref 207. Copyright 2011 Elsevier.

barrier-type oxide films were first formed in sodium arsenate (Na2HAsO4·7H2O) solution, and subsequently the resulting samples were re-anodized in phosphoric acid under electric fields (i) below the threshold electric field (E*) reported by Oh and Thompson for the formation of incipient pores, (ii) close to the threshold field to induce significant anodic oxidation of aluminum, and (iii) well above the threshold field. From microscopic analyses of the arsenic and 18O contents, and the pore morphologies and arsenic distribution in the resulting anodic oxide films, it was found that the field-assisted oxide dissolution is mainly responsible for the formation of incipient pores at oxide formation efficiency (ηj) = ∼20−30%, while field-assisted flow of oxide materials is operative for the growth of major pores at oxide formation efficiency (ηj) = ∼57− 66%.208 The authors suggested that the preferential growth of incipient pores locally increases the current density at the pore bases, which may influence the transport numbers of mobile ions, insertion of electrolyte-derived anionic species into the anodic oxide, and thus oxide viscosity.208 6.3.3. Average Field Model for Steady-State Pore Structure. As discussed in section 6.2.1, after the formation of incipient pores, some large incipient pores develop into major pores by increasing their size through persistent merging with neighboring smaller incipient pores. As the anodization proceeds, the growing major pores readjust their sizes and spatial arrangement to establish equilibrium morphology (i.e., hexagonally close-packed pore distribution). During this period, merging, dying, or even branching of the pores may occur. For a given set of anodization conditions, the barrier layer thickness (tb), pore size (Dp), and interpore distance (Dint, or cell size) in the equilibrium pore structure are mainly determined by the anodizing potential (U). O’Sullivan and Wood suggested a so7510


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called “average field model” to explain the establishment of the equilibrium structure of growing pores.100 The essence of the model is that field-assisted dissolution at the base of growing pores should occur to some different extents until the electric field (E) across the barrier layers eventually becomes the same for every pore, the condition under which pores of an equilibrium size grow at a constant rate, being determined by the average field. The model can be qualitatively explained as follows. Figure 30 schematically shows three adjacent pores

would be enlarged in the lateral direction. Accordingly, for all three cases in Figure 30, rearrangement of pores in conjunction with the self-adjustment of pore size (Dp) and barrier layer thickness (tb) would take place until the electric field (E) across the barrier layers of every pore approaches the average field. 6.3.4. Direct Cation Ejection Mechanism. The average field model discussed above is based on the field-assisted oxide dissolution.100 However, this dissolution mechanism alone does not adequately explain some of the experimental results, in particular, that the quantity of Al3+ ions in the electrolyte after anodization exceeds that of Al3+ ions originating from the formation of the pores. On the basis of pore-filling experiments, Takahashi and Nagayama reported that transport numbers for Al3+ and O2− ions moving across the barrier layer are t+ = 0.4 and t− = 0.6, respectively.46 The result implies that both Al3+ and O2− ions migrate in anodic oxide, with 40% of the ionic current carried by the Al3+ ions and the remainder by the O2− ions, and further that about 40% of the Al3+ ions are lost into the electrolyte without contributing to the porous oxide formation. Considering the porosity (P ≈ 10%) of porous AAO formed in typical anodization conditions,109 30% of Al3+ ions should be lost through a mechanism different from fieldassisted oxide dissolution, even if the latter is still operative. Siejka and Ortega studied the pore formation mechanism by employing 18O tracing techniques.30 They first formed a compact base oxide film on aluminum in electrolyte enriched in H218O and subsequently anodized the sample in H216Oenriched H2SO4 electrolyte for pore formation. From the nuclear microanalyses for 18O-isotope content and depth distribution, the 18O tracer was found to be located at the film surface conserving its initial isotopic concentration (Figure 31b), which was attributed to oxide decomposition inside the

Figure 30. Schematic representation of growing three adjacent pores (a, b, and c) with different pore sizes (Dp1, Dp2, and Dp3) and barrier layer thicknesses (tb1, tb2, and tb3). R1, R2, and R3 are the radii of curvature of the metal/oxide interfaces of the respective pores. ω (∼45.7°) is the angle subtended from the center of curvature to the pore bases. The red lines extending from the metal/oxide interface to the oxide/electrolyte interface represent current lines. The regions marked with “A” correspond to the area of unoxidized metal near the concave ridges.

growing with different pore sizes (Dp1, Dp2, and Dp3) and barrier layer thicknesses (tb1, tb2, and tb3). We assume that pore (b) has the equilibrium dimension of a given set of anodization conditions, in which tb2 and R2 are the equilibrium barrier layer thickness and radius of curvature of the metal/oxide interface, respectively. Although O’Sullivan and Wood assumed that the barrier layer thickness (tb) and the angle (w = 45.7°) in a potentiostatic anodization are constant irrespective of the pore sizes (Dp), we suppose here that only w remains constant while tb changes with Dp, which represents the real cases better. For the three cases given in Figure 30, the electric field (E) across the barrier layers is different, because at a given anodizing potential (U) the field (E) is inversely proportional to the barrier layer thickness (i.e., E = U/tb). Accordingly, fieldassisted oxide dissolution at the pore base would occur to a greater extent for pore (a), as compared to pores (b) and (c), while anodic oxidation of metal would take place to the same extent at the metal/oxide interface due to higher current density (j) according to eq 1, until Dp1, R1, and tb1 have reached their equilibrium values (i.e., Dp2, R2, and tb2). For pore (c), on the other hand, field-assisted oxide dissolution and anodic oxidation reactions would be progressively retarded because of lower electric field (E), until pore (c) has the equilibrium dimension. Unoxidized metal near the hemispherical concave ridges (i.e., the points marked with “A” in Figure 30) is acted upon by two lateral field components of different magnitudes from two neighboring pores. The lateral field component of greater magnitude drives the field-assisted dissolution and anodic oxidation to a greater extent. The corresponding pore

Figure 31. Schematic diagrams illustrating findings of oxygen tracer experiments for the growth of (a) barrier and (b) porous anodic films on aluminum. In each case, the anodizing is carried out first in electrolyte enriched in H218O and second in electrolyte enriched in H216O. Adapted with permission from ref 32. Copyright 2006 The Electrochemical Society.

oxide and reincorporation of the released oxygen to form new oxide at the metal/oxide interface. 18O-tracer studies along with the analyses of the components of ionic current (jtot) have further revealed that the number of oxygen in porous AAO accounts for about 60% of the total ionic current (i.e., jox = 60%·jtot), which is numerically equal to the current efficiency (ηj) of porous AAO formation.29,30 The remainder is associated with the loss of Al3+ ions into the electrolyte (i.e., jloss = 40%· jtot). Siejka and Ortega claimed that pore formation in the absence of oxygen losses should be associated with the direct ejection (jAl = 30%·jtot) of Al3+ ions from the metal/oxide interface into the electrolyte across the barrier oxide and the outward movement (jdec = 10%·jtot) of Al3+ ions produced by 7511


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Figure 32. Cross-section transmission electron micrographs (TEM) of (a) the sputter-deposited aluminum containing a tungsten tracer layer, and following anodization for (b) 180 s, (c) 240 s, and (d) 350 s at 5 mA cm−2 in 0.4 M H3PO4 at 293 K. Reprinted with permission from ref 98. Copyright 2006 Elsevier. The distributions of tungsten tracer in anodic oxide are schematically illustrated above the respective TEM images. (e) Simulated current lines and potential distribution for porous AAO growth in oxalic acid at 36 V (color scale is potential in volts). (f) Simulated flow velocity vectors and mean stress (color scale is dimensionless stress). Panels (e) and (f) reproduced by permission from ref 200. Copyright 2009Macmillan Publishers Ltd.: Nature Materials.

immediately beneath the pores initially lies slightly below the adjacent tracer near cell wall regions (Figure 32b). Upward displacement of the tracer band at the cell wall regions became pronounced upon further anodizations, as compared to the tracer below the pores, where the band is getting fainter because of a reduced tungsten concentration (Figure 32c,d). It was suggested that the fine-tungsten lines along the cell boundaries beneath the tracer layer are associated with the tungsten enrichment in the aluminum adjacent to the aluminum/film interface, because formation of WO3 requires higher Gibbs free energy as compared to Al2O3.98,212,213 The authors pointed out that the observed behavior of the tracer band is contrary to what one would expect from the conventional dissolution-based model of porous AAO growth as follows.32,98 According to the conventional model of pore formation, tungsten tracer should be incorporated into the anodic oxide first at the regions immediately below the pores and then near the cell boundary regions because of the scalloped geometry of the barrier oxide layer. Thus, the tracer band at the pore regions should lie ahead of the tracer at the pore wall regions because of the outward migration of tungsten species, contrary to the experimental observations. If fieldassisted oxide dissolution occurred at the oxide/electrolyte interface, to maintain a constant barrier layer thickness (tb), a tungsten-rich layer with a sharp image contrast (instead of diminished contrast, as in Figure 32c) would have been

oxide decomposition inside the barrier layer toward the electrolyte: jloss = jAl + jdec = 40%·jtot.30 Consistent results have recently been reported by Skeldon and co-workers,161 who observed that the amounts of 18O tracer species in films formed by sequential anodizing in phosphoric acid do not change significantly during the formation of anodic oxide between the barrier and porous stages. The direct cation ejection mechanism is also in line with the recent report by Wu et al., who have experimentally shown that the formation of pores does not occur through oxide dissolution process at the oxide/ electrolyte interface, but through direct ejection of Al3+ ions into the electrolyte.31 However, the process involved in oxide decomposition inside the barrier layer and its role in the dynamic balance between the movement rates of interfaces (i.e., the metal/oxide and the oxide/electrolyte interfaces) need to be specified. 6.3.5. Flow Model for Steady-State Pore Formation. Skeldon et al.32,98,209,210 have investigated the development of pores in porous AAO by performing tracer studies. They anodized Al/W-tracer/Al substrates (W-tracer layer; 3−5 nm thick, Al−30 at. % W, Figure 32a) in H3PO4 solution and investigated the movement of the tracer band (WO3) by TEM and RBS.32,98 During anodic oxidation, tungsten species incorporated into the anodic oxide migrated toward the oxide/electrolyte interfaces at about 0.38 of the rate of Al3+ ions.211 It was observed that the incorporated tracer band 7512


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Skeldon and co-workers213,218 have also investigated the effect of the types of tracer elements on the distributions of tracer layers in anodic alumina. Figure 33 schematically

observed by TEM at the pore bases due to preferential dissolution of Al3+ ions of a higher mobility in anodic oxide. On the basis of the results of tracer experiments, Skeldon and co-workers32,98 proposed the flow mechanism of pore generation, as an alternative process to conventional fieldassisted oxide dissolution. According to the new model, the constant thickness of the barrier layer (tb) during porous AAO formation is maintained by viscous flow of oxide materials from the pore base to the cell boundary. It was suggested by the authors that the displacement of oxide materials is driven by large compressive stresses (∼100 MPa) from electrostriction due to a high electric field (E ≈ 106 V cm−1)97 and also by volume expansion due to oxidation of aluminum,165 and that can be facilitated by the involvement of most of the oxide constituents in ionic transport (i.e., the plastic flow of oxide materials).97,214,215 In most anodizations for porous AAO formation, acid anions from electrolytes are incorporated into the barrier layer through inward migration. They migrate slowly in the barrier oxide, as compared to O2−/OH− ions, due to their relatively large size. As a consequence, the barrier layer exhibits a duplex structure in terms of chemical composition, with an acid anion contaminated outer oxide layer and a relatively pure inner oxide layer adjacent to the metal/oxide interface (see section 5.2).14 The relative thickness of the outer oxide layer to the inner one depends mainly on the nature of the anodizing electrolyte in a given set of anodizations and is invariant throughout anodization. The flow model accounts for incorporation of electrolyte anions into anodic oxide and also their migration behavior in the barrier layer. Incorporated anionic species are transported toward the cell walls in addition to their inward migration in the oxide without dissolutionrelated loss; otherwise, the barrier oxide would eventually contain no incorporated acid anions because of their slow migration. The flow model considers Al3+ ions as the only ionic species being lost into the electrolyte by field-assisted direct ejection.30 Garcia-Vergara et al. have pointed out that bulky acid anions play a key role in the flow of oxide materials, influencing not only the pore and cell dimensions but also the self-organization of pores through the redistribution of stress.32,98,216 The isotopic order and the absence of tracer loss in the earlier 18O tracer studies of Siejka and Ortega (see section 6.3.4)30 have also been recently reinterpreted in terms of the field-induced plastic flow of oxide materials from the pore base toward the cell boundary, during which most of the initially incorporated 18O tracers are displaced to the cell walls.161 The formation of dendritic (or fish-bone-like) pores in a mechanically constrained environment has also been explained in the framework of the flow mechanism.186,217 The flow mechanism of steady-state pore generation has been supported by the theoretical modeling of Houser and Hebert,47,199,200 which highlighted not only the ionic migration under the gradients of mechanical stress and electric potential, but also its implication on the Newtonian viscoelastic flow of oxide materials from the pore base toward the pore bottoms and further into the cell walls (see Figure 32e,f). Yet, the model does not take into account the volume expansion stress at the metal/oxide interface for the viscous flow of the oxide materials. The authors concluded that the compressive stress at the pore base drives the flow of oxide materials in association with the competition of strong anion adsorption with deposition of oxygen ions.200

Figure 33. Schematic diagram comparing (a) hafnium-, (b) neodymium-, and (c) tungsten-containing tracer layers in porous anodic films formed on aluminum in phosphoric acid. Reprinted with permission from ref 213. Copyright 2009 Elsevier.

compares the distributions of hafnium-, neodymium-, and tungsten-containing tracer layers in porous films formed in phosphoric acid. Contrary to what was observed from the tracer studies using tungsten, the hafnium and neodymium exhibited the types of movement behaviors that one might expect from field-assisted oxide dissolution. The authors attributed the observed behaviors of tracer bands to the fast movement rates of the tracer cations (i.e., Hf4+ and Nd3+) in anodic alumina as compared to W6+ ions and also to the loss of fast moving tracer cations into the electrolyte through direct ejection mechanism; Hf4+ ions migrated outward in anodic alumina at about the same rate as Al3+ ions,219 while Nd3+ ions migrated about twice as fast as Al3+ ions (i.e., 6 times faster than W6+ ions).220 Although the authors’ explanation of the unexpected distortion seems to be plausible within the present knowledge of anodization, more systematic experimental investigations using other electrolyte systems and theoretical study are required to fully verify the oxide flow model.

7. SELF-ORDERED POROUS ANODIC ALUMINUM OXIDE (AAO) Studies on porous-type anodization have been mainly led by the surface finishing industry, and hence focused on the development of cost-effective anodizing processes and the improvement of engineering properties of anodized products. Although various anodizing processes have been intensively explored by industry, the size uniformity and spatial ordering of pores have not been considered a major concern. Typical porous AAOs formed by industrial processes, represented by hard anodization (HA), exhibit disordered pore structures with numerous micrometer-sized cracks.221−224 Therefore, classical HA processes have not been implemented in the current nanotechnology research, until the recent development of new HA processes. On the other hand, the mild anodization (MA) 7513


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Figure 34. (Left) A schematic showing a conventional two-step mild anodization (MA) process for self-ordered porous anodic aluminum oxide (AAO): (i) the first long-term anodization, (ii) removal of disordered porous AAO, and (iii) the second anodization at the identical condition to the first one. Representative surface SEM micrographs of the respective samples are shown in panels (a), (b), and (c), respectively. Reprinted with permission from ref 228. Copyright 1998 Springer Science and Business Media. (d) A color-coded SEM image of AAO formed by two-step MA using 0.3 M oxalic acid at 40 V, showing a poly-domain structure. An area with the same color consists of a domain. The pores are color-coded on the basis of the average angle to the six nearest neighbors. Pores that have no apparent hexagonal coordination (i.e., defect pores) are marked with white. Reprinted with permission from ref 226. Copyright 2008 The American Chemical Society.

aluminum substrate maintaining their directional coherency (Figure 34c). Position-defined pore generation during the second anodization has been attributed to the relatively thin native oxide at the bottom of each concave, where the electric field (E) is the highest and the resistance is the lowest.168 In general, porous AAOs formed by two-step MA process exhibit a polydomain structure (Figure 34d).226 The lateral size of the defect-free domain increases with the anodizing time, but is limited to several micrometers. Since the development of the two-step process from oxalic acid-based anodization, a lot of studies have been conducted not only to produce porous AAOs with different pore sizes and densities with an improved arrangement of pores, but also to understand the mechanism responsible for the growth and selforganizing behavior of pores during anodization. Particular efforts have been devoted to exploring the optimum conditions for pore ordering, mainly with sulfuric, oxalic, or phosphoric acid. Studies have indicated that the self-organized growth of ordered pores occurs within a relatively narrow window (known as the “self-ordering regime”) of anodizing conditions (see Figure 36). For a given electrolyte system, in the case of MA, there is an optimum anodizing potential (U) for the best ordering of pores: (i) sulfuric acid (0.3 M H2SO4) at U = 25 V 106,227 for an interpore distance (DMA (ii) oxalic acid int ) = 65 nm; MA (0.3 M H2C2O4) at 40 V for Dint = 103 nm;6,106,168,227,228 (iii) 11 selenic acid (0.3 M H2SeO4) at 48 V for DMA int = 112 nm; and MA (iv) phosphoric acid (0.3 M H3PO4) at 195 V for Dint = 500 nm.109,229 Considerable efforts have been made to explore new self-ordering regimes in a wider range of DMA int . With the idea of controlling self-organization by adjusting the mechanical stress between the metal/oxide interface, Shingubara et al. employed a mixture solution of 0.3 M sulfuric acid and 0.3 M oxalic acid (v/v = 1:1) to change the density of the resulting porous AAO and obtained self-ordered porous AAO with DMA int = 73 nm at U

of aluminum produces self-ordered porous AAOs with uniform pore size (Dp) and interpore distance (Dint), which can be easily tuned by an appropriate selection of anodization conditions. Thus, that process has been intensively utilized in academic research for a wide variety of nanotechnology applications. Yet the MA process is slow (film growth rate = 2−10 μm h−1), and can be conducted within a narrow range of anodizing conditions. In the following sections, we will discuss the conventional MA and the newly developed HA processes, factors governing the structure of porous AAO, and some characteristics of porous AAO that are relevant to nanotechnology applications. In addition, recent attempts to engineer the internal pore structure of AAO by pulse anodization (PA) and cyclic anodization (CA) will be discussed. 7.1. Mild Anodization (MA)

Masuda and Fukuda found that the bottom part of porous AAOs films produced by anodizing of aluminum in 0.3 M oxalic acid at 40 V exhibits a self-ordered pore structure as a result of the gradual rearrangement of the initially disordered pores.5 On the basis of this experimental finding, Masuda and Satoh developed the so-called “two-step anodization” process, by which porous AAOs could be obtained with a highly ordered arrangement of uniform nanopores (Figure 34).6 In a typical two-step anodization, the first anodization process is conducted for more than 24 h. The resulting porous AAO with disordered pores in its top part (Figure 34a) is selectively removed by the so-called “PC-etching” process using an aqueous mixture of 0.5 M H3PO4 and 0.2 M CrO3 at 80 °C.225 The surface of the resulting aluminum is textured with arrays of almost hemispherical concaves (Figure 34b). The second anodization is carried out with the textured aluminum under the same condition employed for the first anodizing. Pores nucleate at the centers of each concave feature and grow normal to the 7514


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= 37 V.230 More recently, Sun and co-workers231 reported an H3PO4-based MA process using aluminum oxalate ((AlC2O4)2C2O4) as an additive to suppress breakdown of porous AAO during anodization at high potential and temperature. They showed that the interpore distance (DMA int ) of porous AAOs can be continuously tuned from 410 to 530 nm by performing anodization with the mixture solution at an anodization potential (U) ranging from 180 to 230 V. Ono et al.17,232 demonstrated the fabrication of porous AAOs with an interpore distance range of 300−600 nm by MA in organic acid electrolytes: DMA int = 300 nm for 5 M malonic acid at 120 V, MA DMA int = 500 nm for 3 M tartaric acid at 195 V, and Dint = 600 nm for 2 M citric acid at 240 V. Yet the spatial ordering and size uniformity of pores of the resulting porous AAOs were by far inferior to those formed by conventional MA processes using sulfuric, oxalic, or phosphoric acid, which would limit the practical applications of such porous AAOs to nanotechnology research. For self-ordered porous AAO formed by MA processes, recent morphological studies indicated that anodizing potential (U) determines the interpore distance DMA int , barrier layer MA 106,109,116 thickness tMA The interpore b , and pore diameter Dp . MA ) and barrier layer thickness (t distance (DMA int b ) increase linearly with anodizing potential (U) with proportionality −1 constants ζMA = 2.5 nm V−1 for DMA int and ARMA = 1.2 nm V 106,109,114 MA for tMA . The observed potential dependence of D b int and tMA is consistent with the earlier reports for disordered b porous AAO.7,100,114 Linear dependence of the pore diameter (DMA p ) on anodizing potential (U) has also been reported both for disordered and for self-ordered porous AAOs. O’Sullivan and Wood100 reported that pore diameter of disordered AAOs increases with anodizing potential at a rate of ζp = 1.29 nm V−1. For self-ordered porous AAO, on the other hand, Nielsch et al.109 proposed that self-ordering of pores requires a porosity (PMA) of about 10% regardless of the specific anodizing −1 conditions by assuming DMA int = 2.5 nm V . Because porosity (PMA) of AAO is given by eq 53, 10% porosity requirement of Nielsch et al. dictates a linear increase of the pore diameter with the anodizing potential at a late of ζp = 0.83 nm V−1, irrespective of anodizing conditions. However, there are a number of recent reports, indicating that the self-ordering of pores can occur at other porosity (PMA) values ranging from 0.8% to 30% depending on the MA conditions, and thus that the anodizing potential (U) is not the only parameter determining the pore diameter (Dp).11,102,232,233 For example, Nishinaga et al. have recently reported fabrication of porous MA AAO with DMA p = 10.4 nm and Dint = 112 nm (porosity, PMA = 0.8% according to eq 53) by H2SeO4-based anodization at 48 V and 0 °C.11 The authors noted that the solubility of anodizing acid electrolyte has important effects on pore diameter and interpore distance, and suggested that the weak solubility of selenic acid under the anodization condition employed caused the formation of 10-nm-scale pores.11 In another example, Chen et al. reported a continuous decrease of the pore diameter, which did not affect either the interpore distance or the barrier layer thickness, when increasing the concentration of polyethylenglycol (PEG) additive in phosphoric acid electrolyte.233 Because the porosity (P) of AAO is determined by the ratio of the pore diameter to the interpore distance (i.e., Dp/ Dint in eq 53), increasing the PEG concentration in the anodizing electrolyte corresponds to decreasing the porosity of AAO. The use of an organic additive to reduce the pore diameter has recently been extended by Martiń et al., who

demonstrated fabrication of porous AAOs with pore diameter less than 15 nm by using an aqueous mixture of sulfuric acid and ethylene glycol (EG).234 Chen et al. noted two possible origins for the decreasing porosity with the addition of organic additive:233 (i) the increase of the effective electric field (E) due to the reduction of the dielectric constant (ε) of the anodizing electrolyte upon addition of an organic modulator, and (ii) the weak chemical dissolution of the pore wall oxide under the protection of organic molecules. Among them, the former is in line with the earlier report by Ono et al. that the Dp/Dint ratio decreases with increasing electric field strength (E) during anodization.232 Su et al. have further explored the effect of electric field (E) on porosity (P). In a series of papers,235−237 they proposed a mechanistic model, proposing that the porosity (P) of AAO is directly governed by the relative rate of water dissociation at the oxide/electrolyte interface, according to the following relation:236 P = 3/(n + 3)

(66)

where n is the amount of water that dissociates per mole of Al2O3 in the oxide dissolution reaction given in reaction 15. The model regards both the dissociation of water and the dissolution of oxide at the oxide/electrolyte interface as important processes, although the latter has recently been largely disputed, as discussed in sections 6.3.4 and 6.3.5. By performing quantum-chemical model computations, the authors showed that the electric field (E) can significantly facilitate heterolytic dissociation of properly oriented water molecules at the oxide surface, that is, field-enhanced water dissociation.41 Under a stronger electric field (E), the dissociation rate of water will be increased (i.e., an increase of the n value in eq 66), and thus the porosity will be reduced. As such, the model appears to adequately explain the dependence of porosity (P) on electric field (E). The above relation was derived from a simple geometric consideration of growing AAO and the mass balance of total oxide-forming oxygen anions produced by the dissociation of water and dissolution of Al2O3. In other words, the original model by Su et al. postulates that all of the oxygen anions produced by water dissociation and oxide dissolution should contribute to the oxide formation at the metal/oxide interface.236 However, by itself this assumption does not fulfill the necessary condition required for the formation of the steady-state pore morphology, that is, the constant barrier layer thickness (tb) maintained by the dynamic balance of the movement rates of metal/oxide and oxide/electrolyte interface, as follows. A close inspection of reaction 15 used in the authors’ model reveals there is a different field dependence for the water dissociation than for the oxide dissolution process. To increase the n value under a stronger electric field (E), water dissociation should occur at a greater rate than oxide dissolution. Under such circumstances, one may expect a progressive increase of the barrier oxide layer thickness (tb) as the anodiziation proceeds. Therefore, other ionic processes may be required to explain the constant thickness of the barrier layer during anodization. It was assumed that Al3+ ions are directly ejected from the metal/ oxide interface without contributing to the oxide formation.235 When MA is conducted outside the self-ordering regime, the degree of spatial ordering of nanopores decreases drastically. For a given acid electrolyte, there is a breakdown potential (UB), above which anodization is accompanied by local thickening, burning, and cracking of the anodic oxide film, caused by a catastrophic flow of current (j) and consequent 7515


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evolution of a large amount of heat:4,18,79,89,94,97,238,239 UB = 27, 50, and 197 V for sulfuric, oxalic, and phosphoric acid, respectively. It has been known for the MA process that anodizations just below the breakdown potentials (UB) yield porous AAOs with the best self-ordered pore structures.18 Ono and co-workers232,240 investigated the self-ordering of nanopores at the local areas of burnt protrusions with a large number of cracks. They found that the locally thickened areas formed by breakdown exhibit domains of highly ordered cell arrangement. In addition, the cell size (Dint) and the barrier layer thickness (tb) of porous AAO at the burnt area were observed to be remarkably smaller than those at the burnt-free area, which is in line with the earlier report by Tu et al.239 These experimental observations imply that current density (j) primarily determines the degree of spatial ordering of pores and the structural parameters (i.e., cell size Dint and the barrier layer thickness tb) of porous AAOs. 7.2. Hard Anodization (HA) Figure 35. (a,b) SEM images showing weak cell junction strength of porous AAOs formed by H2SO4-hard anodization (HA). (c) A schematic illustration showing two different fracture modes of porous AAOs formed by H2SO4-hard (HA) and mild anodization (MA): the A−A′ cleavage plane for HA-AAO and the B−B′ cleavage plane for MA-AAO. Panel (a) reprinted with permission from ref 18. Copyright 2006 The Electrochemical Society. Panel (b) reprinted with permission from ref 113. Copyright 2008 The American Chemical Society.

As mentioned in section 4.1, stable anodization in a given electrolyte is difficult to maintain over breakdown potential (UB) due to the occurrence of burning or breakdown of the growing anodic oxide film. Classical hard anodization (HA) processes adopted in industry have typically been conducted at potentials (U) higher than breakdown value (UB) at the expense of randomly occurring local breakdown of anodic oxide film. The processes take advantage of high speed growth (typically 50−100 μm h−1) of oxide films due to the high current density (j) at an increased anodizing potential (U). However, the resulting porous anodic films are characterized by severely buckled (or uneven) surfaces with numerous cracks and non-uniform (or distorted) pores. Control of the pore size (Dp), interpore distance (Dint), and the aspect ratio of the nanopores is also very difficult. For these reasons, the classical HA processes have not been employed for nanotechnology applications. Various attempts have been made to overcome the problems associated with local breakdown events during HA. The research activity to date includes extending anodizing potential (U) over breakdown potentials (UB) by appropriately tuning the three major pore-forming acid electrolytes (i.e., H2SO4, H2C2O4, and H3PO4), by searching for new anodizing electrolytes, and by efficiently removing the reaction heat. The current density (j) in HA process is typically 1 or 2 orders of magnitude higher than that of conventional MA processes. Thus, Joule’s heating (Q) during HA is 2 or 4 orders of magnitude greater than the ordinary MA processes. The excessive heat during the HA process can not only promote acidic dissolution of the resulting porous AAO, but also trigger local breakdown events.4,239 Chu and co-workers18,110 reported that the breakdown potential (UB = 27 V) in a sulfuric acid-based anodization system can be increased up to 70 V by experimentally aging the electrolyte after a long-term anodizing (i.e., pre-electrolyzing) at 10−20 A hours per liter. Because of the high anodizing potential (U), the current density (j) for stable anodization correspondingly rose significantly, and thus led to a high-speed film growth. The authors were able to fabricate self-ordered porous AAOs with Dint = 90−130 nm at 40−70 V and 160− 200 mA cm−2 at 0.1−10 °C. That condition is far from those of MA (H2SO4: 25 V and 2−15 mA cm−2) but similar to those of classical HA,221,223,241−243 although the authors named this process “high-field anodization”. Porous AAOs formed by this

process exhibited poor mechanical stability, which greatly limits their practical application as templates for various nanofabrications. The produced porous AAOs exhibited weak cell junction strength (Figure 35a), and thus the individual alumina nanotubes could be easily separated upon applying a weak external stress (Figure 35b), unlike the case of porous AAOs formed by MA processes (Figure 35c). Chu et al. pointed out that the concentration of Al3+ ions dissolved in an aged solution plays a key role in the stable growth of porous AAOs at high potentials (U > Ub) and current densities (j), avoiding the burning or breakdown of porous AAO films.18,110 However, the electrochemical action of the aged solution has not yet been clearly understood. In an attempt to determine the effect of aged solutions on breakdown potential (UB), Schwirn et al.113,119 performed a series of HA experiments at the potential range of U = 27−80 V by using sulfuric acid containing different concentrations of Al3+ ions. Their results have shown that stable anodization at high potentials (U > UB) is actually determined by the initial limiting current density (jlimit), not by the solution state. For an oxalic acid-based anodization system, Lee et al.111 have shown that the self-ordering regimes can be extended by performing HA of aluminum. By introducing a thin (ca. 400 nm) porous oxide layer onto an aluminum substrate, gradually increasing the anodizing potential (U) to a target value (100− 160 V) at the rate of 0.5−0.9 V s−1, and effectively removing the reaction heat through direct thermal contact of the aluminum substrate with an underlying cooling plate, the authors could suppress the breakdown of oxide films and grow mechanically stable highly ordered porous AAOs at anodizing potentials of U > 100 V. Their HA process established a new self-ordering regime with widely tunable interpore distances: DHA int = 220−300 nm at U = 110−150 V and j = 30−250 mA cm−2. They suggested that current density j (i.e., the electric field E) is a key parameter governing the self-ordering of pores 7516


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agent for lowering the freezing point of the electrolyte down to −10 °C, but also as a coolant for removing a large amount of reaction heat through its vaporization from the metal/oxide interface. The process allowed the fabrication of self-ordered porous AAOs with various interpore distances: DHA int = 70−140 nm for H2SO4-HA at 30−80 V, 225−450 nm for H2C2O4-HA at 100−180 V, and 320−380 nm for H3PO4-HA at 195 V.118,122 Microscopic investigations of HA-AAOs have shown that the barrier layer thickness (tb) increases at a rate of ARHA = 0.6−1.0 nm V−1 with respect to anodizing potential (U) depending on the current density (j),18,111,113,118,119 which is smaller than ARMA ≈ 1.2 nm V−1 for MA-AAOs.7,100,114 Lee et al.111 attributed the reduced ARHA (i.e., anodizing ratio) to the high current density (j) involved in the HA process in accordance with the high field conductivity theory, which predicts an inversely proportional dependence of tb to the logarithm of current density (j) at a given anodization potential (U) (see eq 1). Figure 36 summarizes the self-ordering potentials (U) and corresponding interpore distances (Dint) of porous AAOs formed by MA and HA in three major pore-forming electrolytes (i.e., H2SO4, H2C2O4, and H3PO4). It was found that HA-AAOs exhibit a reduced potential (U) dependence of the interpore distance (Dint) with a proportionality constant ζHA = 1.8−2.1 nm V−1,16,110,111,113,118,122,244,245 as compared to self-ordered MA-AAOs (i.e., ζMA = 2.5 nm V−1).109 On the other hand, HA experiments performed at potentiostatic conditions (i.e., U = constant) have also shown that the interpore distance (DHA int ) of HA-AAOs decreases with the current density (j), indicating that anodizing potential (U) is not the only parameter determining the cell size of porous AAOs.16,111,113,118 Lee et al. assumed that high mechanical stress at the metal/oxide interface due to high current density j (i.e., high electric field strength, E) may be responsible for the reduced ζHA.111 In other words, energetically unfavorable fieldinduced mechanical stress (i.e., electrostriction pressure) was suggested to be accommodated by reducing the cell size to increase the surface area of the metal/oxide interface. This is reminiscent of the decrease of critical wavelength (λc) of surface undulations upon increase in the internal stress (σ) for an initially flat surface (i.e., λc → ∞ as σ → 0) according to the model of Srolovitz:246,247

Figure 36. (a−f) Representative SEM images of MA- and HA-AAOs. Panel (a) reproduced with permission from ref 226. Copyright 2008 The American Chemical Society. Panel (b) reproduced with permission from ref 106. Copyright 1998 AIP Publishing LLC. Panel (c) reproduced with permission from ref 229. Copyright 1998 The Japan Society of Applied Physics. Panel (d) reprinted with permission from ref 110. Copyright 2005 Wiley-VCH Verlag & Co. KGaA, Weinheim. Panel (e) reproduced with permission from ref 111. Copyright 2006 Nature Publishing Group. Panel (f) reproduced with permission from ref 118. Copyright 2006 IOP Publishing. Reprinted with permission from IOP Publishing. (g) Self-ordering regimes in MA (filled symbols) and HA (open symbols) by using H2SO4 (black symbols), H2C2O4 (red symbols), H2SeO4 (green symbol), and H3PO4 (blue symbols). The black solid lines represent the linear regressions of the data with correlation parameters of ζMA = 2.5 nm V−1 and ζHA = 1.8−2.0 nm V−1. Data for H3PO4-HA (△) show current density (j) dependence of the interpore distance (Dint) at a fixed anodizing potential (U = 195 V). Data derived from refs 106,227 for H2SO4-MA, refs 5,106,168 for H2C2O4-MA, ref 11 for H2SeO4, refs 109,229 for H3PO4-MA, ref 113 for H2SO4-HA, ref 111 for H2C2O4HA, and ref 118 for H3PO4-HA.

λc = πMγ /σ 2

(67)

where γ is the surface energy and M is the biaxial modulus of a solid. On the other hand, on the basis of the experimentally observed tensile to compressive stress transition during galvanostatic anodization in 0.4 M H3PO4, Proost et al. have reported that the factor controlling the interpore distance (Dint) is not likely to be an internal stress-induced surface perturbation, but rather an electrostatic energy-induced surface instability with the critical perturbation wavelength (λc) given by the following equation:175

at a given anodizing potential (U), which is in line with the suggestion by Ono et al.232,240 The authors also found that HA process produces porous AAOs at 25−35 times faster growth rates, as compared to conventional MA processes. Interestingly, the porosity (PHA) of HA-AAOs was found to be 3.3−3.4%, which is about one-third of the porosity value (PMA ≈ 10%) that was proposed as a requirement for self-ordered MA-AAOs by Nielsch et al.109 The experimental method has also been applied to sulfuric and malonic acid-based HAs.16,113 As mentioned above, HA of aluminum is accompanied by a large amount of reaction heat. Accordingly, the reaction heat should be properly removed for stable anodization at high potential (U > UB) and current density (j). To this end, Li et al.118,122 added ethanol (C2H5OH) to the aqueous anodizing electrolytes (e.g., H2SO4, H2C2O4, and H3PO4) and conducted stable HA at high potentials (U) and current densities (j). In their HA processes, the added ethanol served not only as an

λc = (4πtbγ /ε0εrE2)1/2

(68)

where tb is the barrier layer thickness, ε0 and εr are the vacuum and the relative permittivities of the oxide, respectively, and E is the electric field.248 The instability equation predicts that the interpore distance (Dint) is a function of (tb/E2)1/2 with the slope of 2π/(γ/ε0εr)1/2. Interpore distances measured during the growth of porous anodic alumina and titania agree closely with those predicted by their electrostatic energy-based perturbation criterion.175 7517


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Figure 37. (a) Current (j)−time (t) transient showing spontaneous current (j) oscillation during hard anodization of aluminum at 200 V with an unstirred 0.3 M oxalic acid. An instant solution agitation was introduced at the time indicated by the arrow in (a). (b) An enlarged j−t curve after the electrolyte agitation, together with the corresponding pore structure of AAO. Reprinted with permission from ref 103. Copyright 2010 Wiley-VCH Verlag GmbH & Co. KGaA, Weinheim.

Recently, Lee et al.103 accidently found that porous AAOs that experienced spontaneous current oscillations (amplitude ∼800 mA cm−2) during a potentiostatic HA under unstirred electrolyte conditions exhibit modulated pore structures. From anodizing experiments performed at potential (U) = 140−200 V, the authors observed that peak profiles of the oscillating currents are symmetrically sinusoidal with relatively small but increasing amplitudes at the early stage, while asymmetric with larger and uniform amplitudes at the lager stage of anodization (Figure 37a). The oscillation period was observed to increase with time. Microscopic investigation of the resulting porous AAOs revealed that the pattern of pore modulations matches exactly with the detailed profile of oscillating current peaks (Figure 37b); the pore diameter and the segment length of modulated pores increase with the amplitude and the period of oscillating currents, respectively. The authors attributed the spontaneous current oscillations to the periodic concentration change of anodizing electrolyte at the pore bottom under unstirred electrolyte (i.e., a diffusion-controlled anodic oxidation of aluminum). On the basis of the observed dependence of the pore diameter on the anodizing current density, they suggested that the internal pore structures of porous AAOs can be tailored by judiciously controlling the

current density (j) during anodization of aluminum under potentiostatic conditions. 7.3. Pulse Anodization (PA)

Ordered porous AAOs with tailor-made internal pore structures may provide a new degree of freedom in templated syntheses of novel nanomaterials,249−251 and also serve as ideal platforms for investigating the adsorption and separation behaviors of diverse particles, ions, or biologically important molecules.252−254 On the basis of the experimental finding that HA of aluminum produces porous AAOs with one-third lower porosity (P) than MA (i.e., PHA ≈ 3% for HA and PMA ≈ 10% for MA), Lee and co-workers111 have demonstrated fabrication of AAOs having periodically modulated nanopore diameters along the pore axes by combining MA and HA processes (Figure 38). In their method, each step for modulation of pore diameter (Dp) required a tedious manual exchange of the anodizing electrolytes to satisfy both MA and HA conditions. In an attempt to avoid this problem, they have recently developed an approach for continuous modulations of internal pore structure of porous AAOs by pulse anodization (PA) of aluminum under a potentiostatic condition using sulfuric or oxalic acid solution.112 Historically, pulse anodization (PA) of aluminum was developed in the early 1960s.255,256 The process has 7518


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Figure 38. SEM image of porous AAO with modulated pore diameters prepared by cycling mild anodization (MA) and hard anodization (HA). A magnified image of the area marked with a white rectangle is shown as an inset. Reproduced with permission from ref 111. Copyright 2006 Nature Publishing Group.

popularly been employed in the aluminum industry to produce anodic films of high technical quality (i.e., improved corrosion and abrasion resistance) at an efficient rate of production.257−259 However, the process has been out of view in academic research in the past four decades, and has not been applied to the development of nanostructured materials due to the non-uniform and disordered pore structures of the resulting anodic alumina. In a newly developed PA process, a low potential (UMA) and a high potential (UHA) are alternately pulsed to achieve MA and HA conditions, respectively. A typical pulse profile is schematically shown in Figure 39. A current recovery behavior is observed for MA pulses, similarly to the current evolution at the early stage of ordinary MA processes (Figure 23a). Current is high at the initial stage, drops to a minimum value, and then gradually increases to reach a steady value after passing an overshoot.112,260,261 On the other hand, upon applying an HA pulse, the current density (j) increases steeply (typically, jHA ≈ 10−102 × jMA) for a short period of time and then decreases exponentially, which is the typical anodization kinetics of HA of aluminum under a continuous potentiostatic condition.111 In the anodization of aluminum, Joule’s heat (Q) for a given period of time (t) is proportional to the square of the current density (j):

Q = Ujt = R bj 2 t

Figure 39. Schematics showing potential (U)−current (j) relation during potentiostatic pulse anodization (PA).

of the Joule’s heat. The heat generated by an HA-pulsing can be effectively dispersed during the subsequent MA-pulsing.224,263 In a typical PA process, the current density (j) changes periodically to the values determined by pulsed potentials (i.e., jMA for UMA and jHA for UHA; jMA < jHA). As a consequence, the resulting porous AAO exhibits a lamellar structure with alternate stacking of MA-AAO slab with a smaller pore diameter and HA-AAO slab with a larger pore diameter (Figure 40a−c), where the thickness of each oxide layer is determined by the pulse durations (i.e., τMA for MA-pulse and τHA for HA-pulse).112,224 In addition to such a structural modulation, porous AAOs formed by PA also exhibit a periodic

(69)

where Rb is the resistance of the barrier layer.262 HA of aluminum produces a large amount of reaction heat due to the high anodic current density (j). The generated heat would increase the proton activity of the acid electrolyte, and thus may cause undesired acidic dissolution of pore wall oxide or even local burning of the growing porous AAO via catastrophic local flow of current (i.e., electrolytic breakdown). Therefore, the heat should be properly removed during anodization at high current density (j). PA provides an effective way of dissipation 7519


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Figure 40. (a) Scheme for the preparation of porous AAOs with modulated pore diameters by pulse anodization (PA). (b) False-colored SEM image of AAO formed by H2SO4 PA. AAO slabs formed by MA and HA pulses are indicated by MA-AAO and HA-AAO, respectively. (c) SEM image showing modulated pore diameter. (d) SEM image of 3D stacks of MA-AAO slabs, obtained by selective removal of HA-AAO slabs by chemical etching. Reproduced with permission from ref 112. Copyright 2008 Nature Publishing Group.

ultrasonic treatment to separated individual nanotubes from the sample. As was discussed in section 7.2, MA and HA exhibit different dependences of the barrier layer thickness (tb) on anodizing potential (U); ARMA ≈ 1.2 nm V−1 for MA and ARHA = 0.6−1.0 nm V−1 for HA. According to eq 1, the anodizing current density (j) is exponentially proportional to the inverse of the barrier layer thickness (tb). In PA of aluminum, therefore, when the anodizing potential is changed from a higher UHA to a lower UMA, the current density drops abruptly from jHA to a minimum value and then increases gradually to a value (jMA) determined by UMA (i.e., current recovery). It was reported that the time required for a complete recovery of anodizing current depends on the chemical nature of the barrier oxide (i.e., the content of anionic impurities), the electrolyte temperature, and the potential difference between UHA and UMA.112,264 For PA in three major pore-forming acid electrolytes, Lee and Kim noted that the time required for current recovery increases in the order: H2SO4 < H2C2O4 < H3PO4.264 Further, they pointed out that PA of aluminum in H2C2O4 or H3PO4 electrolyte is difficult to achieve continuously within a reasonable period of time because of the retarded current recovery, especially at a low temperature and a large potential difference.264 To resolve the problem associated with the slow current recovery, they increased gradually anodizing potential prior to pulsing a high anodizing potential.264 Porous AAOs with tailor-made internal pore geometries could be conveniently prepared by employing potential pulses with deliberately designed periods and

compositional modulation along the pore axes. TEM-EDX point analysis of porous AAO formed by sulfuric acid-based PA revealed that the amount of anionic impurities (mostly SO2− 4 ) in HA-AAO slabs is about 88% higher than that in MA-AAO slabs, which was attributed to the high current density (jHA) during HA-pulsing.112 As discussed in section 5.2, anodic oxide contaminated with anionic impurities exhibits a poor chemical stability against an oxide etchant (e.g., 5 wt % H3PO4). By taking advantage of higher level of anionic impurities in HAAAO slabs, Lee et al.112 could completely separate the MAAAO slabs from a single as-prepared porous AAO by performing selective chemical etching of HA-AAO slabs (Figure 40d). Structural modulation of porous AAOs can also be achived by galvanostatic PA, where current pulses satisfying MA and HA conditions are periodically applied. As discussed in section 7.2, H2SO4-AAO formed under HA conditions shows fairly weak junction strength between cells.110,112,113 On the basis of this experimental finding, Lee and co-workers263 explored a convenient route for the mass preparation of uniform alumina nanotubes with prescribed lengths. The authors applied periodically galvanic MA and HA pulses to achieve continuous modulations of pore diameter and also to weaken the junction strength between cells. Alumina nanotubes, the length of which is determined by the HA-pulse duration (τHA), could be obtained by immersing the resulting porous AAO into an aqueous mixture of 0.2 M CuCl2 and 6.1 M HCl, followed by 7520


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Figure 41. Schematics showing (a) the experimental process for the fabrication of AAO with tailor-made pore structures by pulse anodization (PA) and (b) a generalized form of a potential pulse employed in pulse anodizations. Uj and τij define the repeating unit of potential waves, where Uj = the potential at the time tj with U1 = U5 (j = 1−4), τij = tj+1 − tj, i = the pulse number (i = 1, 2, 3...). Representative SEM image of porous AAOs with modulated pores prepared by pulse anodization; (c) τ11 = 36 s, (d) τ11 = 144 s, (e) τ11 = τ21 = 36 s, τ31 = τ41 = 144 s, and (f) τ11 = τ21 = 144 s, τ31 = 36 s, τ41 = 144 s, τ51 = τ61 = 36 s. Other parameters were fixed at U1 = 80 V, U2 = 140 V, U3 = U4 = 160 V, τi2 = τi4 = 0 s, τi3 = 0.2 s. The repeating units of pulses are shown as insets in the respective images. Reprinted with permission from ref 264. Copyright 2010 IOP Publishing.

beginning with first cycle having a gradually increasing amplitude of current, then a double-profiled cycle, and last a series of triangular galvanic cycles.265 AAOs with periodically perforated pores (i.e., pores with nanoholes along horizontal directions) were also fabricated by chemical etching of cyclic anodized porous AAOs.266

amplitudes (Figure 41). This capability for engineering internal pore geometry may provide a unique opportunity in templated synthesis of nanowires and nanotubes with modulated diameters, and also in utilization of porous AAOs with 3D periodic pore structures in photonic applications. 7.4. Cyclic Anodization (CA)

7.5. Anodization of Thin Aluminum Films Deposited on Substrates

The concept of structural engineering of porous AAO by combining MA and HA has been put forward by Losic et al.,265,266 who developed a new anodizing process, called cyclic anodization (CA). They employed periodically oscillating current signals with different cyclic parameters (i.e., period, amplitude, and profile) during anodization of aluminum to achieve structural modulations in porous AAOs (Figure 42).265 Porous AAOs with modulated pores of different shapes (circular- or ratchet-type), diameters, and lengths were prepared by applying current signals of deliberately chosen cyclic parameters. Microscopic investigation of AAOs formed by CA processes indicated that the structural details of pores follow exactly the applied current profiles. The authors pointed out that the transitional anodization (TA) mode, the transition from MA to HA condition (Figure 42d), is important for structural engineering of pores in CA process. It was proposed that the TA mode creates a pore, whose internal geometry is directed by the characteristics (i.e., profile, period, and amplitude) of the applied cyclic signal. They further showed nanostructuring of porous AAOs with distinctive, hierarchical internal pore structures by employing multiprofiled current signals: for example, three different successive CA steps,

In addition to overcoming their brittle characteristics, AAOs grown with thin Al films deposited on substrates of choice would potentially offer much broader application than those on bulk aluminum foils. The substrate could be insulators (e.g., glass),267−270 semiconductors (Si and TiN),271−277 non-valvemetal (e.g., Cu, Ag, Au, Pt, etc.)270,278−283 or valve-metal (e.g., Ti, W, Nb, Zr, Ta, etc.)-coated Si substrates,284−290 and transparent indium thin oxide (ITO).291−298 AAOs formed by anodizing thin Al films on substrates have been utilized not only as patterning masks, but also as templates for fabricating various functional nanostructures, including arrays of Si nanoholes,271 carbon nanotubes (CNT),287 magnetic nanowires,289,293 thermoelectric nanowires,270 photocatalytic nanowires or nanotubes,293,296 and valve-metal oxide nanodots or nanorods.277,284,286,288 Previous works have indicated that the anodizing behavior of the interfacial area of the Al and substrate material depends on the underlying material in terms of anodization kinetics and structure of the barrier layer (Figure 43). When anodizing Al films on insulating substrates (e.g., SiO2) in three major pore7521


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Figure 42. (a) Schematics of the galvanostatic cyclic anodization (CA) for structural modulation of porous AAO. (b,c) SEM images of AAOs with modulated pores fabricated by asymmetrical current signals (i.e., exponential saw-tooth) with two different amplitudes. (d) Influence of current amplitude on the pore shape and the length of modulated pore segments. Anodization modes (i.e., MA, TA, HA) associated with corresponding anodization currents are marked on the pore structures and graphs. Reprinted with permission from ref 265. Copyright 2009 Wiley-VCH Verlag GmbH & Co. KGaA, Weinheim.

Figure 43. Schematic current (j)−time (t) curves during anodization of thin Al films on (a) insulating substrates, (b) non-valve-metal-coated (or ITO) substrates, and (c) valve-metal-coated substrates. (d−f) Schematic cross-sections of the respective porous AAOs. The solid curve in (b) corresponds to j−t curve during anodization of silicon substrate without conducting surface layer.

Nb, Zr, Ta, etc.) are filled with corresponding oxide nanodots or nanorods, protruding from the barrier layer (Figure 43c,f).284−290 Among the above-mentioned substrate materials, the anodization of Al film on a Si substrate has been the most extensively investigated. When anodizing an Al/Si substrate, after the complete anodization of the Al film, oxidation of the Si surface forms SiO2 nanodots under an inverted barrier oxide layer. The thickness of the inverted barrier layer is significantly reduced as compared to that formed on bulk Al foils.271,272,299,300 Seo et al. suggested that formation of interfacial voids and the inversion of the barrier layer has a mechanical origin, involving multiple process stages, as follows

forming electrolytes (i.e., H2SO4, H2C2O4, and H3PO4, Figure 43a,d), the completion of the anodization is marked by a sharp decrease in current density (j) and color change.267−270 The barrier layer has a U-shaped morphology like that of AAO formed from bulk Al foils. Al just underneath the cell boundary remains unoxidized in the form of discrete nanometer-sized particles, which are trapped between the alumina barrier layer and the insulating substrate. Unlike bulk Al foils, however, the barrier layers of AAOs formed by anodization of Al films on conductor-coated substrates (e.g., Si, Pt, Au, ITO, etc.) are characterized by inverted morphology with interfacial voids (Figure 43b,e).270−274,276−282,291−297 On the other hand, pores of AAOs formed on valve-metal-coated substrates (e.g., Ti, W, 7522


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Figure 44. Cross-sectional SEM images (top) and a schematic (bottom) showing the process of void formation: (a) touching of the barrier layer with Si surface, (b) flattening of pore bottoms due to stress accumulation, (c) void nucleation to minimize the stress, and (d) formation of inverted barrier layer. SEM images were adapted with permission from ref 276. Copyright 2007 Elsevier.

(Figure 44).276 When the barrier layer touches the Si surface, the residual Al underneath the cell boundary region is laterally confined and thus cannot accommodate the volume expansion stress due to the rigid Si substrate without interfacial restructuring to create the necessary additional space (i.e., void). The driving force for void nucleation is the stress pushing the Si substrate downward. On the other hand, upward stresses are exerted on the bottom of each pore, resulting in the inversion of the barrier layer curvature (i.e., detachment of the barrier layer). The electric field for anodizing the remaining Al is locally concentrated on the pore edges. As a result, a dendritic branching occurs at the edge of each pore bottom. The upward stress increases due to void growth during anodizing of the residual Al. The inversion behavior allows the barrier layer to be farther away from the Si substrate, which relieves the stress burden while the barrier layer becomes thinner upon further anodizing.276 From TEM nanoprobe energy dispersive X-ray spectroscopy (EDS), Seo et al.276 found that the inverted barrier layer surrounding a void has a compositional bilayer structure with a thin Al-rich region near the void and relatively thick Al-poor region near the pore bottom, the origin of which is likely to be associated with the detachment process of the barrier layer during void formation. Anodizing Al film/Si substrate for an extended period of time results in a local oxidation of Si by electrolyte infiltrating through the channels between the pore bottoms and voids, forming SiO2 nanodots just underneath the voids,276,300 accompanied by a violent evolution of oxygen gas bubbles. As a result of the gas evolution, the porous AAO typically delaminates from the Si substrate. Similar phenomena have also been observed for anodization of Al films on ITO/glass substrates, although Chu et al.291,292,295,298 have successfully anodized thin Al films on ITO/glass substrate. This may be ascribed to improved physical bonding of the AAO as a result of the sputter deposition of highly energetic Al atoms on the ITO/glass substrate. In the case of Al films on non-valve-metalcoated substrates (e.g., Cu, Pd−Au, Ag, Pt), prolonged anodization results in detrimental breakdown or dissolution of metal during the anodization.270 Accordingly, anodization for a prescribed time is required for each substrate. For electronic applications, it is desirable to grow a barrierlayer-free AAO directly on conductor-coated substrates. Sander and Tan demonstrated the fabrication of a barrier-layer-free AAO membrane by anodizing Al film deposited on an Aucoated substrate, followed by chemical etching of the barrier layer in 5 wt % H3PO4.278 Yet their chemical etching process resulted in enlargement of the pore diameter, due to the

isotropic nature of anodic alumina etching. Moreover, several minutes of anodization of the sample resulted in detachment of the AAO from the substrate surface. Yang et al. anodized an Al/ Au/Si substrate in an effort to circumvent the need to open the barrier layer.281 Yet, the Al/Au bilayer system formed Au−Al intermetallic phases, which catalyzed a deleterious oxygen evolution reaction, causing detachment of the AAO membrane from the substrate.281,283 For anodization of Al/Pt/Si substrates, a barrier layer thinning process based on stepwise voltage reduction279 or the use of a reverse bias in KOH270 have been employed to minimize pore widening. However, the stepwise voltage reduction process resulted in the bifurcation of pores near the barrier layer. Under the reverse polarization conditions, the Pt underlayer catalyzed the electrolysis of water, violently evolving H2 gas and thus delaminating porous AAO from the substrate. To improve physical bonding between the porous AAO and a conductor-coated substrate, a thin interlayer (typically, ca. 5-nm-thick Ti) between the Al and the conductor layer has been introduced.282,283,301 Yasui et al. introduced a thin layer of Ti (1.5 nm) between the Al and Pt layer to serve both as an adhesion promoter and as a barrier eliminating thin TiO2 in 5 wt % H3PO4.282 Oh and Thompson have successfully demonstrated a selective barrier perforation process by anodizing Al films on a W(60 nm)/Ti(15 nm)-coated silicon substrate.289,290 Their process is based on selective dissolution of the anodized valve-metal oxide (i.e., WO3) in a pH 7 phosphate buffer solution. During removal of the WO3, isotropic pore wall etching did not take place due to the neutral nature of the etching solution. After the selective removal of WO3 from the pore base, Oh and Thompson could directly grow Ni or Pt nanowires by electrodeposition into the pores of AAO, in which the exposed base metal layers served as cathode.289,290 Pringle215 first theoretically analyzed the anodizing behaviors of superimposed valve-metal layers (i.e., bilayered metal films). He predicted that the metal order would be conserved if the anodic oxide of the superimposed metal was less resistive, while the metal order would be partially inverted through finger penetration of oxide by the underlying anodic oxide if the superimposed anodic oxide was more resistive. He also pointed out that these phenomena may be changed by the effects of transport number, relative rate of cation migrations, oxide structure, and PBR.215 Later, this model was experimentally confirmed by Shimizu and co-workers,214,302−304 who observed that the random penetrations of lanceolate oxide fingers from the less resistive underlying oxide layer into the more resistive upper oxide layer take place. Recent works by Mozalev and co7523


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Figure 45. SEM images of anodic tantala nanorods formed by multi-step anodization of Al/Ta/Si substrate: (a) before and (b) after removal of AAO. (c−e) Schematic cross-sectional views of (c) the Al/Ta bilayer anodized at 53 V in 0.2 M H2C2O4, (d) after potentiostatic re-anodizing from 53 to 310 V in 0.5 M H3BO3, and (e) after pore widening followed by re-anodizing from 53 to 310 V in 0.5 M H3BO3. Reprinted with permission from ref 307. Copyright 2004 The Electrochemical Society.

stamp (mold) onto the aluminum by mechanical pressure (i.e., nanoimprinting), followed by anodization. The SiC imprint stamps were fabricated with an array of convex features of desired arrangements in a limited dimension by electron beam lithography (EBL) technique.308−311 Each shallow indentation formed on the aluminum surface defined the position of pore growth by initiating pore nucleation at the initial stage of anodization, and thus led to a perfect arrangement of pores within the patterned area (Figure 46a). Masuda et al. further extended the method to fabricate pore array architectures with square- or triangle-shaped pore openings in square or triangular arrangements (Figure 46c,d).309 In a pre-pattern on aluminum, the missing sites of the pattern can be compensated automatically during anodization, if the distance between the missing site and the nearest patterned sites satisfies the potential (U)−interpore distance (Dint) relation required for the self-ordering of pores (i.e., ζMA = Dint/U = 2.5 nm V−1, section 7.1).311 The size of the pores formed at those missing sites is smaller than that of pores formed at the patterned sites (Figure 46e,f). By utilizing this self-compensation ability in AAO growth, Smith and co-workers have demonstrated the fabrication of porous AAOs with hybrid circular-diamond and circular-triangular-diamond pore cross-sections.312 By anodizing aluminum with surface pre-patterns in nonequilibrium

workers have indicated that in certain electrolytes, anodizing of thin Al deposited on a layer of valve-metals (e.g., W, Ti, Ta, Nb) results in the filling of AAO pores with nanodots or nanorods of the corresponding metal oxides (Figure 45).20,286,305−307 More recently, Chu et al.288 obtained similar experimental results from anodization of Al film on a Zr-coated glass substrate. Mozalev et al.307 pointed out that the filling of AAO pores with other valve-metal oxides is possible due to the higher PBR and cation transport number (t+), as compared to those of Al; for example, the PBR for Ta/Ta2O5 = 2.5 and t+Ta = 0.24−0.5.

8. LONG-RANGE ORDERED POROUS AAO Porous AAOs formed under the self-ordering conditions exhibit a poly-domain structure, where each domain contains hexagonally ordered nanopores of an identical orientation and is separated by the boundaries. The domain boundaries are characterized by defect pores (white dots in Figure 34d). For most nanotechnology applications, porous AAOs with uniform pore size and long-range ordering of pores are required. Masuda et al.308 first reported fabrication of ideally ordered porous AAOs with a single-domain configuration over a few mm2 area. The process involved pre-patterning of the aluminum surface by transferring the pattern of a hard SiC 7524


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Figure 46. (a) Schematics showing process for ideally ordered porous AAO. (b−d) SEM micrographs of porous AAOs with different hole array architectures: (b) circular, (c) square, and (d) triangular pore openings. (e−j) SEM micrographs showing self-compensated pore formation; (e−g) surface prepatterns with missing sites of pattern and (h−j) the respective porous AAOs formed by anodization. The black arrows in (h) indicate pores formed at the missing sites of pattern in (e). SEM micrographs shown in (f,g,i,j) highlight formation of porous AAOs with (f,i) hybrid circulardiamond and (g,j) hybrid cicular-diamond-triangle pore cross-sections; the scale bars are 500 nm. Panel (b) was reprinted with permission from ref 308. Copyright 1997 AIP Publishing LLC. Panels (c,d) were reprinted with permission from ref 309. Copyright 2001 Wiley-VCH Verlag GmbH & Co. KGaA, Weinheim. Panels (e,h) were reprinted with permission from ref 311. Copyright 2001 AIP Publishing LLC. Panels (f,g,i,j) were reprinted with permission from ref 312. Copyright 2008 AIP Publishing LLC.

cracking the underlying brittle substrate, high mechanical pressure also causes structural damage of the imprint stamp after several times of pattern transfer. In attempts to resolve the above-mentioned problems, other pre-patterning techniques have been proposed: evaporation of thin aluminum film on ordered arrays of self-assembled Fe2O3 nanoparticles (NPs) on silicon followed by mechanical striping,318 resist-assisted or direct patterning by focused-ionbeam (FIB) lithography,319,328−330 and direct nanoindentation using a tip of a scanning probe microscope (SPM).320,331 Yet these methods are not practical for pre-patterning over an extended area due to the limited ordered area (ca. 2 μm2) of nanoparticles for the first technique, and due to the serial nature of the patterning process for the other two techniques. For large-scale fabrications of long-range ordered porous AAOs on fragile substrates, pre-patterning of thin aluminum films by holographic lithography,321 block-copolymer lithography,322 and nanosphere lithography (NSL)323 has been developed. Fabrication of single domain porous AAOs over 2-in. silicon wafer has been also realized by directly anodizing thin aluminum films deposited on a lithographically generated SiO2 mask, in which ordered SiO2 holes underneath the aluminum film define the position of pore growth by effectively guiding the electric field during anodization.332 Most of these techniques have demonstrated their effectiveness in fabrications of ideally ordered porous AAOs with tunable interpore distance (Dint) on brittle substrates over an extended dimension. However, they are inherently incapable of generating multiple copies of surface pre-patterns on aluminum. Soft lithography utilizing elastomeric poly-dimethylsiloxane (PDMS) stamp or nanoimprint lithography (NIL) on polymeric resist has been demonstrated as a versatile method for multiple transfers of a master pattern onto various substrates.325,333,334 Pattern transfer in these lithographic techniques can be achieved at pressures 103−105 times less than those used in hard stamping directly onto aluminum. Lai

tessellation arrangement (Figure 46f,g), the authors found that the cell geometries in the resulting porous AAOs are determined by the arrangements of the unpatterned and patterned pore sites and also direct the cross-sectional shape of pores: circular, diamond, and triangular pores, respectively, in regular, elongated, and partially compressed hexagonal cells (Figure 46i,j). The authors attributed the evolution of diamond and triangular pore cross-sections to a coupled effect between the thick pore wall oxide and the longer segment length of cellboundary bands (section 5.2), suppressing or eliminating the influence of the smaller cell segments.312 Porous AAOs with sharp-featured non-circular pore cross-sections can be used as templates in synthesizing functional nanostructures for enhanced sensing in localized surface plasmon resonance (LSPR) and surface-enhanced Raman spectroscopy (SERS).313 Stimulated by the works of Masuda and co-workers, several groups have developed various surface pre-patterning methods to fabricate single-domain AAOs with tunable interpore distances, Dint (Table 2). Earlier research in this direction was mostly devoted to the development of an economic way of producing hard imprint stamps with large lateral dimensions. Surface pre-patterning by mechanical nanoindentations using optical diffraction grating,314 Si3N4 mold fabricated by deep-UV lithography,315 wafer-scale Ni mold fabricated by laser interference lithography (LIL),317 and self-assembled monoor multi-layer of nanospheres316,327 is effective in terms of process cost and pre-patterning area. However, these approaches have an intrinsic limitation, in that they rely on the transfer of stamp patterns onto the aluminum surface under high mechanical pressure (typically, 5−100 kN cm−2) and therefore are not suited for pre-patterning of thin aluminum layers deposited on technologically more relevant substrates for device integration, such as fragile silicon or glass. In general, as the pattern density on an imprint stamp increases, the higher the required applied pressure must be in this pattern transfer protocol. In addition to severely deforming the aluminum and 7525


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a

100

∼10 cm area

wafer scale

O

O

100

277

2 cm × 2 cm

O 2

84

>1 mm2 area

O

45

∼300

1 cm × 1 cm

O

O

100

1 μm × 1 μm

100

63 481 500 150 200 13

patterned area

3 mm × 3 mm 5 mm × 5 mm 4-in. wafer >cm2 area wafer scale 2 μm × 2 μm

O

O

X X X X X X

reported min Dint (nm)

Ni mold was replicated from a master pattern fabricated by EBL, and its pattern was transferred onto a thermoplastic resist; subsequently, the resulting resist pattern was transferred onto an Al surface by Ar-ion beam milling; imprint pressure: 0.5 kN cm−2 ordered arrays of nanoindents were generated by wet-etching of Al through pre-patterned polymer masks that were created by SFIL on Al surface; imprint pressure: 10 or at a high temperature, either porous AAO dissolves and/or the pores of the AAO template become closed due to the higher rate of metal deposition relative to the rate of mass transfer of Au(I) and formaldehyde down to the pores. Yu et al.394 reported that Au nanotubes formed by electroless deposition are characterized by nanoclustered morphology, in which the size of the gold nanoclusters increases with the pH of the plating bath. They showed that the size of the gold nanoclusters determines the catalytic properties of Au nanotubes embedded in AAO membrane. Further, they showed that Au nanotubes/ AAO composite membranes can be reused in the catalytic conversion of 4-nitrophenol into 4-aminophenol in the presence of NaBH4 as a reductant. Activation of the porous AAO membranes can also be accomplished by immobilizing PdNPs on the pore wall surfaces.395−397 Fabrications of arrays of Co,398 Ni,398 Cu,398,399 Pd,400 Pt,400 binary398 or ternary401 metal alloy nanotubes, nanowires,402 and nanocones403 have been demonstrated by electroless deposition of metals into PdNPimmobilized porous AAO membranes.

Figure 53. SEM images of multisegmented metal nanotubes with a stacking configuration of Au/Ni/Au/Ni/Au along the nanotube axis: (a) false-colored cross-sectional SEM image of as-prepared metal nanotube-AAO composite. (b,c) SEM images of mutisegmented metal nanotubes after removal of the AAO template. Adapted with permission from ref 375. Copyright 2005 Wiley-VCH Verlag GmbH & Co. KGaA, Weinheim.

9.2. Electroless Deposition (ELD)

Electroless deposition of metals into porous templates was pioneered by Martin and co-workers.378−381 Most of the earlier works were devoted to electroless deposition of gold within the pores of track-etched polycarbonate membranes. The typical deposition process is composed of the following three general steps:378,382,383 (1) Sensitization: This is accomplished by immersing the porous polymeric template into an aqueous solution of SnCl2 to deposit Sn2+ onto the surfaces of the membrane. (2) Activation: This is accomplished by dipping the Sn2+sensitized polymeric membrane into a AgNO3 solution, which yields metallic AgNPs on the membrane surfaces through a surface redox reaction (see reaction 70). (3) Electroless deposition: This entails immersing the activated polymeric membrane into a gold plating solution, in which the surface-bound AgNPs act as catalyst for the reduction of Au+ to yield AuNPs on the membrane surface through the following reaction. Au+aq + Ag 0surface → Ag 0surface + Ag +aq

9.3. Sol−Gel Deposition

Sol−gel processing has developed into a versatile protocol for the stoichiometric synthesis of diverse nanocrystalline materials. In general, sol−gel chemistry involves the hydrolysis of precursor molecules under an acidic condition to prepare a suspension of colloidal particles (i.e., sol) and the subsequent condensation of the sol particles to obtain a gel. The resulting gel is then calcined to obtain the desired material. Metal alkoxides in organic media or inorganic salts in aqueous media can be used as precursors. Martin and co-workers pioneered the sol−gel porous AAO templating method for the synthesis of various semiconductor or insulator oxide nanostructures (nanowires and nanotubes) including TiO2, ZnO, WO3, V2O5, MnO2, Co3O4, and SiO2.404−406 Since then, many investigators have employed the sol−gel deposition method to prepare a vast variety of oxide nanostructures to investigate, for example, photovoltaic (TiO2),407,408 gas sensing (SnO2, ZnO),409,410 ferromagnetic (CoFe2O4),411,412 ferroelectric (BiFeO3, SrBi2Ta2O9, PbZr0.52Ti0.48O3),413−415 superconducting (YBa2Cu3O7-δ, Bi2Sr2CaCu2Oy),416,417 and luminescence (Y2O3:Eu, TiO2, ZnO:Dy) properties.418−420 Typically, when depositing sol−gel into oxide nanopores, an AAO is directly dipped into a solution containing sol particles for a given period of time. After thermal treatment, either nanowires or nanotubes of inorganic oxide are formed within

(71)

The surface-bound AuNPs act as autocatalysts for the reduction of Au+ to metallic Au0 in the presence of a reducing agent (e.g., formaldehyde). Au deposition starts at the entire surface of membrane (i.e., pore walls and membrane faces). As a result, the surfaces of pore walls and the faces of the membrane can be coated with a thin Au layer, yielding Au nanotubes within the pores and continuous Au films on the faces of the membrane. The resulting gold nanotube/polymer composite membranes have been successfully used for diverse applications, for example, for selective ion-transport,379,382−387 (bio)molecule separations,380,388,389 biosensing, and electroanalysis.390−393 7531


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the authors were able to synthesize very thin single-crystalline TiO2 nanowires (diameter ≈ 10 nm) by performing heattreatment and subsequently removing the AAO template. As a surface modification protocol, the sol−gel template technique offers a large degree of freedom in nanotechnology applications of porous AAO. Lee et al.425 demonstrated that sol−gel-derived silica nanotubes/AAO composite can be used as a synthetic bio-nanotube membrane for separating two enantiomers of a chiral drug, in which individual silica nanotubes act as nanometer-sized chromatography columns in parallel. They deposited thin-walled (

Porous Anodic Aluminum Oxide: Anodization and Templated

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