Journal of The Electrochemical Society, 152 �12� F203-F212 �2005� F203

Proton Diffusion and Electrochromism in Hydrated NiOyand Ni1−xVxOy Thin FilmsE. Avendaño,a,b,d,z A. Azens,c G. A. Niklasson,a and C. G. Granqvista,*aDepartment of Engineering Sciences, The Ångström Laboratory, Uppsala University,SE-751 21 Uppsala, SwedenbCenter of Science and Engineering of Materials, CICIMA, Universidad de Costa Rica, San José,Costa RicacChromoGenics Sweden AB, Uppsala Science Park, SE-751 83 Uppsala, Sweden

Electrochromic hydrated nickel oxide and nickel vanadium oxide thin films were deposited by reactive dc magnetron sputtering.

Optical modulation was effected by insertion and extraction of protons. The proton diffusion coefficient D̃ during the insertion/

extraction process was determined by galvanostatic intermittent titration. We analyzed D̃ as a function of changes in the stoichi-ometry and found deep minima at certain H/Ni ratios. These minima were interpreted, within the bleached state, as due to astructural phase transition from �-nickel hydroxide to �-nickel hydroxide. The optical absorption increased rapidly at the H/Niratio corresponding to the second phase transition. At the transition to the colored state, our data were consistent with a structuralphase transition from �-nickel hydroxide to �-nickel oxy-hydroxide, in accordance with the Bode reaction scheme. Those struc-tural changes were corroborated by infrared reflection-absorption spectroscopy.© 2005 The Electrochemical Society. �DOI: 10.1149/1.2077308� All rights reserved.

Manuscript submitted March 8, 2005; revised manuscript received July 22, 2005. Available electronically October 24, 2005.

0013-4651/2005/152�12�/F203/10/$7.00 © The Electrochemical Society, Inc.

Electrochromic materials change their optical appearance revers-ibly between transparent and dark upon charge insertion/extraction.1

Materials that color upon insertion are called cathodic, while mate-rials that color upon extraction are called anodic. Thin films of hy-drated nickel oxide possess anodic electrochromic properties, as dis-covered by Svensson and Granqvist in 1986.2

The nickel oxide changes its color from transparent to brownupon charge extraction. In electrochromic devices combining mate-rials based on WO3 and NiO, the brown color of nickel oxide iscomplementary to the blue color of tungsten oxide, both togetheryielding a neutral �gray� appearance in the dark state.1,3-5 Applica-tion areas for electrochromic devices include architectural “smartwindows” backed by glass or polyester, variable-transmittance eye-wear, variable-reflectance mirrors, information displays, and vari-able emittance surfaces.1,3,4 Nickel oxide and hydroxide have beenwidely studied in the context of battery electrodes as well.

Electrochromic films can be produced by a variety of techniques,with reactive dc magnetron sputtering from metallic targets being ofparticular importance for large-scale, large-area manufacturing.1,3,4

The magnetic character of Ni is an obstacle for this technique,though, which justifies an in-depth analysis of additives that renderNi nonmagnetic without compromising the superior electrochromicproperties on Ni-based films to any significant extent. Vanadium is acommon additive to remove the magnetic property of Ni-based sput-ter targets, and therefore we included nickel vanadium oxide in thepresent investigations.

There has been a long-time controversy in the literature concern-ing the type of ions transferred in nickel oxide upon coloration, withevidence being presented for both proton2,6-9 and OH− transport.10

Proton diffusion coefficients of the order of 10−12 cm2/s have beenreported.11,12 A change in the hydrogen concentration upon colora-tion was detected by nuclear reaction analysis.2,13 Covering the coat-ing with a palladium layer and cycling in a KOH electrolyte pro-vided additional evidence of proton transfer, as apparent fromreflectance modulation observed from the back side of the system.7

Because the layer of palladium catalyzes hydrogen and no otherspecies can penetrate the Pd film, the change of absorption wasattributed to the insertion/extraction of protons. A correlation wasfound between a high concentration of OH− species, a low resistiv-ity of the oxide film, and an increase of the optical absorption.10

* Electrochemical Society Active Member.d Present address: Brazilian Synchrotron Light Laboratory, LNLS, CEP 13084-971,

Campinas, SP, Brazil.z E-mail: Esteban.Avendano@Angstrom.uu.se

Many reactions have been proposed in order to explain themechanism by which the change in color occurs. The Bode reactionscheme is frequently used; it was introduced to account for thecharging/discharging phenomena in Ni electrodes using KOH aselectrolyte14,15 and is expressed by

�-Ni�OH�2 ↔ �-NiOOH

� ��-Ni�OH�2 ↔ �-NiOOH

�1�

Several alternative reactions have been proposed in the literature.Thus, charging/discharging phenomena, assuming interchange ofprotons, were presented as the topochemical mechanism9

�-Ni�OH�2 ↔ �-NiOOH2−z + zH+ + ze− �2�

or, in a more simple way, as1

Ni�OH�2 ↔ NiOOH + H+ + e− �3�

Another mechanism was proposed, assuming the interchange of hy-droxyl groups, according to16

Ni�OH�2 + OH− ↔ NiOOH + H2O + e− �4�

for nickel hydroxide and

NiO + OH− ↔ NiOOH + e− �5�

for nickel oxide.17 Finally we mention, for completeness, an expla-nation of the coloration mechanism by the creation of superoxideaccording to18

3Ni�OH�2 + M+ ↔ M�NiO2�3 + 6H+ + 5e− �6�

where M+ is the cation of the supporting electrolyte. From this briefsurvey it is evident that more work is needed to clarify the colora-tion mechanism, and this paper has such clarification as one of itsgoals.

The present work deals with sputter-deposited nickel-oxide-based films made under varied conditions. Nearly metallic filmswere obtained at low oxygen content in the sputter plasma, whileoptimum electrochromic performance was found for films depositedat intermediate oxygen and hydrogen content. Those films havegood crystalline structure and high porosity. We note that earlierwork has recognized that porosity and small grain size, leading to alarge surface-to-bulk ratio for the grains, increases the electrochro-mic activity.19-21 Finally, an overdose of either oxygen or hydrogengave films with reduced charge capacity and limited optical modu-lation range, presumably connected with the increase of density andlowered crystallinity. Optimization of the deposition process was

F204 Journal of The Electrochemical Society, 152 �12� F203-F212 �2005�F204

reported elsewhere.21 The role of vanadium addition to the nickeloxide, in order to facilitate the film manufacturing by sputtering,was a small one, essentially limited to giving some additional opti-cal absorption in the bleached state.

The Galvanostatic intermittent titration technique,22 known asGITT, was used to analyze the behavior of the chemical diffusioncoefficient as a function of changes in the stoichiometry. This led tonew results on the mechanism of coloration. At least two phasetransitions are involved: one change, presumably from �-nickel hy-droxide to �-nickel hydroxide, within the bleached state, and an-other one to the colored state between �-nickel hydroxide and�-nickel oxyhydroxide. These assignments follow the Bode reactionscheme.

Experimental

Nickel-oxide-based thin films were made by reactive dc magne-tron sputtering using a versatile deposition system23 based on aBalzers UTT 400 unit. The sputter targets were 5-cm-diam plates ofnonmagnetic NiV0.08 with 99.95% purity. Sputtering was conductedin a plasma of Ar + O2 + H2. The gases were 99.998% pure. Thetotal sputter pressure was 30 mTorr, and the sputtering power was200 W. Depositions took place onto substrates positioned 13 cmfrom the targets. The substrates used for optical and electrochemicalmeasurements were glass plates precoated with a layer of indium tinoxide �ITO� �i.e., In2O3:Sn24�, with a resistance/square of 60 �.Graphite substrates were employed for compositional determina-tions using Rutherford backscattering spectrometry �RBS�. Deposi-tion rate was obtained from sputter time and ensuing film thicknessrecorded by surface profilometry across a step made in the film.Typical deposition rates and film thicknesses were 0.4 nm/s and200 nm, respectively.

Elemental composition was determined by RBS. Ions of 4He at2 MeV, with a scattering angle of 166.1°, were used to determinethe concentration of the heavy atoms in the film. Structural analysiswas performed by X-ray diffraction �XRD� using a Siemens D5000diffractometer operating with a grazing incidence angle of 1° inparallel-beam geometry with twice the diffraction angle 2� lyingbetween 30° and 120° and a wavelength of 1.540598 Å, correspond-ing to Cu K� emission.

Cyclic voltammograms were taken in a three-electrode cell con-sisting of a working electrode, which was the hydrous nickel-oxide-based film deposited onto a glass/ITO substrate, a counter electrodeof Pt, and a standard reference electrode of Ag/AgCl �KCl withconcentration 1 M�. The voltammograms were recorded in a 1 MKOH solution, as in some prior work on NiO,2,25 and the experimentwas carried out with an Autolab PGSTAT10 potentiostat.

Optical properties in the UV-visible �UV-vis� to near-infrared�NIR� range were measured on a Perkin-Elmer double-beam spec-trophotometer. Specifically transmittance �T� and reflectance �R�were determined in the 300 � � � 2500 nm wavelength range atnormal or near-normal incidence, using an integrating sphere. Back-ground corrections were taken for each scan. Barium sulfate wasused as a reference for the reflectance measurements. In situ cyclicvoltammetry, together with transmittance measurements, were ac-complished at 350 � � � 800 nm using an Ocean Optics instru-ment.

Infrared reflectance was determined using a Perkin-Elmerdouble-beam spectrophotometer �model 983�. Data were taken inthe 200-4000 cm−1 wavenumber range with p polarized light and anangle of incidence of 60°. A thin film of gold was used as reference.

GITT was carried out with the Autolab PGSTAT10 potentionstat.In this technique one applies a current step for a short period of timeso that intercalation or deintercalation of ions takes place from theelectrolyte, and subsequently the circuit is opened and the cell al-

lowed to relax.22 The chemical diffusion coefficient D̃ can be deter-mined through analysis of the intercalation or deintercalation stepand the relaxation mode. All measurements were carried out in a1 M KOH solution using a three-electrode cell.

Physical Characterization

Deposition regions.— Three distinct deposition regions—denoted 1, 2, and 3—can be identified when sputter deposition ofnickel-oxide-based films is carried out under reactive conditions in aplasma of Ar with varying amounts of O2.21,26 When this content isincreased, the films evolve from nearly metallic �region 1� to trans-parent �region 2�, and then from transparent to brown �region 3�.Films in region 1 do not exhibit electrochromism, films belonging toregion 2 have a very high activity and good charge capacity, andfilms in region 3 have low activity and poor optical properties. Thedeposition rate drops precipitously at the transition between regions2 and 3. Normal deposition rates were found to lie between50 nm/min for a partly oxidized target and 10 nm/min in the over-oxidized mode of the target. These characteristics are similar for allnickel-oxide-based films. The present paper deals only with filmsproduced by sputtering from NiV0.08 and Ni targets, but many othermaterials have been reported in Ref. 26 and 27.

Composition.— Elemental compositions were determined byRBS. Defining the composition as Ni1−xVxOy, the three regions re-ferred to above correspond to x � 0.11 and y � 0.71 for region 1,0.05 � x � 0.10 and 1.45 � y � 1.75 for region 2, and x � 0.11and y � 2.07 for region 3. As already pointed out, films with supe-rior electrochromism were found in region 2. Our results are consis-tent with those of an independent ion-beam analysis.13 An increaseof the O2 flow in the gas mixture increased the oxygen content in thefilms. Furthermore, an addition of H2 to the gas mixture producedfilms with improved crystallinity and also increased their chargecapacity and transparency.

Density.— The density lay between 3.6 and 4.2 g/cm3 in theoptimized electrochromic films obtained with intermediate O2 andH2 flows during the sputtering. At the transition between regions 1and 2, the density was approximately 6.6 g/cm3, and at the transi-tion to the overoxidized region 3 it was approximately 4.9 g/cm3.Similar values of the density were found for nickel oxide films madeby electron-beam evaporation.1,28,29

Structure.— XRD was used to study the structure of the as-deposited films. Specifically, bleached and colored films, made sothat they belonged to region 2 and deposited onto ITO-coated glass,exhibited cubic nickel oxide �Bunsenite� patterns for 30° � 2�� 120°, which is in agreement with previous work onNiO.1,7,19,25,30-33 No evidence of phases containing hydrogen wasfound in any of the films by using conventional XRD equipment.This at first sight surprising result can be understood if only theoutermost part of the grains that were too thin to be detectable34 takepart in the electrochromic coloration.

Effective grain size L and strain function S were determinedfrom the full width at half maximum fwhm of the XRD peaks usingthe Scherrer equation, which accounts for the grain size broadening�g and the lattice strain broadening �s produced by dislocations anddefects.35 The pertinent relationship is

fwhm�2�� = �g + �s =0.9 �

L cos���+ S tan��� �7�

i.e.,

cos��� fwhm�2�� =0.9 �

L+ S sin��� �8�

where � = 1.540598 Å and � is the Bragg angle in radians. Thefwhm was obtained by fitting each peak to a Gaussian curve. Plot-ting the fwhm times the cosine vs the sine of the Bragg angle, theintersection of the curve yields the broadening contribution of theeffective grain size and the slope yields the broadening produced bythe strain in the film. The strain was calculated assuming an isotro-pic variation along all directions.

Figure 1 shows that the effective grain size lies between 10 and25 nm depending on the H content during the thin-film deposition.

2

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This parameter should be understood as the equivalent diameter ofthe average volume in the film that diffracts coherently. The strain,which is also plotted in Fig. 1, is related to distortions caused bydislocations, defects, and vacancies. Small grain size, implying alarge surface-to-bulk ratio, may be beneficial for the electrochromicactivity, as pointed out before.19-21

Earlier work of ours36 used extended X-ray absorption fine struc-ture spectroscopy on films similar to the present ones and demon-strated that the vanadium atoms substitute nickel in a NiO-typestructure, i.e., nickel and vanadium appeared to form a mixed-oxidephase in the film.

Electrochemical and Optical Characterization

Cyclic voltammetry.— Cyclic voltammograms �CVs�, i.e., mea-surements of current density j vs voltage V, were taken for filmsbelonging to regions 2 and 3. Figure 2 shows CV data pertaining tofilms made with different gas mixtures used during the depositions.The CV for a relative oxygen flow of 6% is representative of theoveroxidized region. We found large differences between films pro-duced under conditions pertaining to different regions. The electro-chromism of films belonging to region 3 is characterized by lowcharge capacity and small optical modulation. Differences of around50%–90% in the charge capacity were seen for those latter filmswith regard to films made under optimum deposition conditions.Furthermore, films of region 3 displayed severe aging effects al-ready after 20–30 cycles.

A stabilization of the electrochromic films, prepared under favor-able conditions, took place after 10–20 cycles depending of thedeposition conditions. Figure 3 presents the evolution of the CVfrom the as-deposited state to the stable state for one of the films ofregion 2, specifically deposited at O2 = 1.92% and H2 = 1.92%. Theincrease in the intensities of the cathodic and anodic peaks impliesthat there is a growth of the charge capacity of the film during thefirst 10–20 cycles. This may be interpreted as a transformation ofthe nickel oxide to nickel hydroxide, as we discuss later. The shiftsof the peaks, indicated by straight dotted arrows, suggest that thereis poor electronic conductivity through the films.

The same film as the one in Fig. 3 was analyzed after stabiliza-tion. Data are reported in Fig. 4 for different values of the voltagesweep rate. Scrutinizing the changes of the CV when this rate wasvaried, we found that �i� the potentials for the different peaks shiftaway from the standard electrode potential, indicating that there theelectronic conductivity through the film is poor,37 �ii� independentlyof the cycle number, the ratio between the intensities of the cathodicand anodic peak is less than unity, �iii� the intensities of the cathodicand anodic peaks increase with an increase of the sweep rate, and�iv� the shoulder at the coloration potential around 0.4 V disappearsat very slow sweep rate. These results are consistent with the reac-tion mechanism that involves a charge transfer of the type

Figure 1. Effective grain size and strain function for Ni1−xVxOy films, de-posited with sputter parameters belonging to region 2, vs the hydrogen frac-tion in the plasma. The O2/Ar ratio was kept constant at 2%.

Ni+2 ↔ Ni+3. It consists of two chemical steps, the first one beingreversible and the following one irreversible. The latter follows fromthe fact that the shoulder for voltages higher than 0.4 V disappearsat very slow sweep rates.37 Potentials higher than 0.43 V led todecomposition of the electrolyte. The irreversible process could beresponsible for the degradation of the films in this electrolyte, butstress and fatigue due to the insertion of charge also need to be takeninto account. We found degradation in the edges of the samples,corresponding to approximately 5% of their surface, after more than120 cycles in 1 M KOH for all films.

Diffusion coefficient: Randles-Sevčik equation.— This equationcan describe the behavior of the anodic current peak density jP ataround 0.3 V in the CV as a function of the sweep rate at which

the potential is scanned, the chemical diffusion coefficient D̃, thetemperature T0, the diffusion length d, and the maximum intercala-tion charge qm. Those variables are connected by the short timedomain of finite space diffusion Randles-Sevčik equation accordingto38

Figure 2. CVs for Ni1−xVxOy films deposited with the shown fractions ofoxygen and hydrogen in the sputter plasma. An example of films depositedwith parameters belonging to region 3 is also shown �dashed curve�. A scanspeed of 10 mV s−1 was applied in a solution of 1 M KOH. Arrows show thesweep direction.

Figure 3. Evolution of CVs from the as-deposited state to charge stabiliza-tion for a Ni1−xVxOy film prepared with the shown fractions of oxygen andhydrogen in the sputter plasma and cycled in 1 M KOH at a sweep rate of10 mV/s. Solid arrows show the scan direction and dotted arrows indicatethe evolution of the peaks.

F206 Journal of The Electrochemical Society, 152 �12� F203-F212 �2005�F206

jP = �0.199F

RgT0

qm

d D̃1/2 1/2 �9�

where F and Rg are Faraday’s constant and the gas constant, respec-tively.

Figure 5 shows a Randles-Sevčik plot for the current density �inmA/cm2� of the anodic peaks extracted from Fig. 4. We found goodlinearity between the peak intensity and the square root of the sweeprate. This linear dependence, marked with a dashed straight line,represents the range of rates for which the process is controlled bydiffusion phenomena. The intercept of the straight line close to zerocurrent provides clear evidence that in the experiment we reduce themass transfer by migration and reduce the faradaic losses in theprocess. The slope of the linear part of the curve indicates that thediffusion coefficient of the extracted species is 2.5 10−16 cm2/s.

In order to estimate how much charge can be extracted from anickel vanadium oxide film, we used the value of the charge capac-

Figure 4. CVs at six different sweep rates for a Ni1−xVxOy film in 1 MKOH. The film was deposited with oxygen and hydrogen ratios of 1.92% inthe sputter plasma. Arrows show the sweep direction.

Figure 5. Randles-Sevčik plot for the anodic peak current � jP� vs square rootof the scan rate �� for a Ni1−xVxOy film cycled in 1 M KOH. Data weretaken from the curves in Fig. 4.

ity for the lowest sweep rate and calculated the number of ionsequivalent to this charge. The sweep rate for this film was0.05 mV/s, and the ratio of monovalent ions to nickel atoms wasapproximately 0.60. A similar result has been reported previouslyusing nuclear reaction analysis.13

Electrochromism: relation between charge and opticaldensity.— In order to illustrate the optical response upon insertion/extraction of charge, we used the film for which data were presentedin Fig. 4. The sweep rate was 0.1 mV/s. We recorded the transmis-sion at � = 550 nm and the charge simultaneously. The optical den-sity �OD� was calculated using Beer’s law by1

T��� = e−OD �10�

The reflectance changes are small from the bleached to the coloredstate, and hence one can use the magnitude of the OD as an approxi-mation for the absorption.

The specific charge, defined as Q/Qcol, was obtained by dividingthe charge Q at each potential with the maximum coloration chargeQcol, in this case being the total charge of the cathodic bleaching.The extra charge due to the O2 evolution in the electrolyte was nottaken into account. The results are summarized in Fig. 6. At around−0.55 V, we have a small absorption in the visible due to the ex-traction of charge and/or residual color due to the vanadium presentin the films.34 This appears as an OD of about 0.1 for voltagesaround zero. The maximum absorption, as found from the OD, oc-curs at around 0.42 V. The application of high positive voltages,above 0.55 V, does not contribute to the electrochromic process andonly accelerates the oxygen evolution in the electrolyte.

Electrochromism: coloration efficiency.— Optical spectra of thestudied films, deposited under the conditions of region 2, werenickel-oxide-like.1,33,39 Transmittance and reflectance were recordedin the as-deposited, colored, and bleached states. The main differ-ences between these states are most prominent in T��� within thevisible region, and only minor changes can be observed in R���. Thepotentials at which the spectra were recorded correspond to the onesof maximum bleaching and coloring �i.e., when the current densityis equal to zero�. Combining the results of the CV with the opticalmeasurements, we can determine the spectral coloration efficiency�CE� by the relation40

Figure 6. Specific charge �Q/Qmax� and OD at the wavelength � vs voltagefor a Ni1−xVxOy film. The film was deposited with the shown fractions ofoxygen and hydrogen in the sputter plasma and was cycled in 1 M KOH at0.10 mV/s.

F207Journal of The Electrochemical Society, 152 �12� F203-F212 �2005� F207

CE =

ln� �1 − RColored�2 TBleached

�1 − RBleached�2 TColored�

�QExchanged charge�11�

Results are shown in Fig. 7. The main changes take place in the UVand visible, whereas the changes are small in the NIR. Three mainpeaks can be distinguished in Fig. 7: two prominent ones are locatedat wavelengths 340 and 445 nm, and a third broad peak can bediscerned at around 600 nm. This last peak is confirmed in Fig. 9. Ingeneral, CE is high compared to values reported in otherwork.1,29,30,39 The influence of the sputtering conditions on the CE isstrong at short wavelengths and weak at long wavelengths.

The CEs for 300 � � � 500 nm correlate with the crystallinityof the films, and films with high CE showed large intensities of theirXRD peaks. We could see direct evidence connecting the colorationefficiency with the grain size, presented in Fig. 1, and with the initialcontent of oxygen in the film.

Before commenting on the results stated above, it is appropriateto clarify two main issues: First, the results presented here deal withthe mechanism of coloration prior to stabilization in sputter-deposited optimized thin films. We note that during stabilizationthere are two reactions occurring in parallel, one related to transferof hydrogen and another one to irreversible uptake of OH− groups inthe outermost parts of the material. The initial oxygen content in theas-deposited films enhances the latter reaction; clear evidence of thisfact was found from X-ray photoelectron spectroscopy �XPS� dataobtained by synchrotron radiation.34 It is outside the scope of thispaper to discuss the initial evolution of the films. However, we em-phasize that major problems for explaining the mechanism of col-oration in nickel-oxide-based materials ensue from three factors: �i�the material that has been produced is not always well defined, �ii�stabilization of the material may not be ensured, and �iii� the appliedpotential may have been too high so that irreversible insertion ofpotassium ions �for the case that KOH is used as electrolyte� hastaken place. A second issue that needs comment is that much of thework published before has dealt with films belonging to what werefer to as region 3. In this case the low crystallinity and the highoxygen content lead to severe degradation of the material prior tostabilization; such films are difficult to analyze because they areunstable and, furthermore, in most cases useless for practical deviceapplication.

Considering the discussion the previous paragraph, we concludethat films with small grain sizes and with initially low oxygen con-tents had low CEs as a consequence of a small amount of activematerial on their surfaces, whereas films with more oxygen andlarger grain sizes showed enhanced CEs because the loss of active

Figure 7. CE vs wavelength � for Ni1−xVxOy films prepared with the shownfractions of oxygen and hydrogen in the sputter plasma. Also shown areO/Ni ratios determined by RBS.

surface area is compensated with an excess of oxygen. Finally, filmswith small grain sizes and high oxygen contents exhibit high CEsbecause of the large surface area with high amount of active mate-rial. This behavior supports the notion that the large inner surfacearea of the films is connected to the electrochromic activity. Thecoloration process takes place in the outermost parts of the grainsbecause it depends on the content of oxygen that is located there. Itshould be noted that the films, irrespective of their state of colora-tion, exhibit a crystalline NiO-type structure, for which the excess ofoxygen can only be accomodated on the surface. Finally, a high CEwas observed for films with a large charge capacity, as measured bycyclic voltammetry.

Electrochromism: proton transport.— With the object of show-ing that proton transfer is responsible for the optical modulation, weused a palladium coating as a membrane that only allows H+ topenetrate through the film. In this way, it is possible to filter out allother ions, such as OH− or K+ present in the electrolyte, becausethey are too large to pass through the Pd coating. We did not see any

Figure 8. CVs for Ni1−xVxOy films prepared with the shown fractions ofoxygen and hydrogen in the sputter plasma. Data are given for a film coatedwith 25 nm of palladium �solid curve� and without such a coating �dashedcurve�. A scan speed of 10 mV s−1 was applied in a solution of 1 M KOH.Arrows show the sweep direction.

Figure 9. Spectral coloration efficiency �CE� for Ni1−xVxOy films preparedwith the shown fractions of oxygen and hydrogen in the sputter plasma andalso reported in Fig. 8. The curve for the film coated with a 25-nm-thicklayer of palladium �solid line� is multiplied by a factor of two. A scan speedof 10 mV s−1 was applied in a solution of 1 M KOH.

F208 Journal of The Electrochemical Society, 152 �12� F203-F212 �2005�F208

activity in a 25-nm-thick layer Pd alone in a study using a 1 M KOHsolution. Some previous work7 also employed a Pd film to single outthe effect of H+ transport.

A 200-nm-thick nickel vanadium oxide film on an ITO-coveredglass substrate was coated with a 25-nm-thick layer of Pd. Figure 8shows the CV for this working electrode in an electrochemical cellwith three electrodes as described before. The CV of a similar filmthat was not coated with Pd is shown for comparison. One can seethat the shape of the CV is similar for the coated and the uncoatedoxide film, except for the extra activity at the negative potentials.The peak at −0.474 V in the cathodic sweep for the coated film canbe due to hydrogen evolution; alternatively, it may be related to theshoulder at 0.033 V in the anodic direction. This shoulder is in-creased because the Pd coating promotes hydrogen dissociation. It isimportant to notice that the value of the coloration potential does notchange �being approximately 0.42 V vs Ag/AgCl�, implying that theabsorbing compound is the same for both films. The bleaching po-tential shifts from −0.42 V to − 0.14 V vs Ag/AgCl.

CE, calculated from Eq. 11 without correction for the reflectance,is plotted in Fig. 9 for a film of nickel vanadium oxide with andwithout a Pd coating. The wavelengths where the CE shows peaksare the same in the two cases. The similarities in the main featuresof the CE in Pd-coated and uncoated films suggest that the modula-tion of the optical absorption in the film is due exclusively to theinsertion/extraction of hydrogen. Optical modulation has been re-ported earlier for nickel oxide films coated with a 100-nm-thick Pdfilm, in which case the reflectance was recorded from the back sideof the sample.7

Infrared reflection-absorption spectroscopy.— This technique iswell established for studying adsorbed substances or thin films ofinsulator or semiconductor materials on metal surfaces.41 Disregard-ing the transmittance, the absorptance can be expressed as

A = 1 − Rp �12�

where Rp is the reflectance of p polarized light. We now assume that�i� the magnitude of the dielectric function for the metal substrate �inthis case ITO� is always larger than that of the absorbing layer, �ii�the metal is an isotropic medium, �iii� the metal has a negative

Figure 10. Energy loss function obtained from measurements with p polar-ized light at 60° angle of incidence in the wavenumber region between 3850and 2950 cm−1. Three different voltages were applied corresponding to thefully bleached �−0.40 V�, intermediate �0.20 V�, and colored state �0.45 V�.

dielectric constant in the infrared region, and �iv� the metal is aperfect conductor. Then, after solving Maxwell’s equations with theappropriate boundary conditions and carrying out tedious algebra,Eq. 12 takes the form

1 − Rp = 4�

cd Im� − 1

����

sin2���cos���

�13�

with d being the film thickness, � the incidence angle, c the velocityof light in vacuum, and � the angular frequency. In the analysisbelow we used p polarized reflectance, measured at an incidenceangle of 60° in the infrared region, to calculate the energy lossfunction from Eq. 13.

A 200-nm-thick film of nickel vanadium oxide deposited onto anITO-coated glass substrate was investigated after stabilization bycycling 70 times in a 1 M KOH solution. The measurements werecarried out ex situ after application of three different potentials.Referring to the voltammogram in Fig. 3, this corresponds to thebleached state �−0.40 V�, an intermediate state �+0.20 V�, and thecolored state �+0.45 V�. Before the measurements, the film wasrinsed in distilled water.

Figure 10 shows results for the three potentials in the wavenum-ber region between 3850 and 2950 cm−1. The broad band at around3400 cm−1 is indicative of water absorption.42,43 The stretching bandof nickel hydroxide and oxyhydroxide, denoted �O–H��, lies closeto the edge at around 3650 cm−1.43-45 This edge undergoes a smallred-shift �cf. the dashed lines� as the applied potential is increased,implying that the strength of the bond weakness and that the bondlength between oxygen and hydrogen increases. The region between2300 and 1200 cm−1 is shown in Fig. 11. The features in the datacorrespond to deformed water, denoted �HOH��, and carbonateimpurities.42 Finally, Fig. 12 shows the energy loss function at lowfrequency, specifically for the region between 860 and 360 cm−1.The discussion to follow is focused on two vibration bands: first, thebroad libration band �O–H�l

44 �defined as an oscillation in the ap-parent aspect of a secondary body �i.e., O–H� as seen from the

Figure 11. Energy loss function obtained from measurements with p polar-ized light at 60° angle of incidence in the wavenumber region between 2300and 1200 cm−1. Three different voltages were applied corresponding to thefully bleached �−0.40 V�, intermediate �0.20 V�, and colored state �0.45 V�.

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primary object �i.e., Ni� around which it revolves� in the regionbetween 560 and 400 cm−1, and second, the stretching band�Ni–O�� at around 590 cm−1.46,47

Chemical Diffusion Coefficients Determined by GITT

Overview.— Transport properties were investigated by GITT. Westudied a series of films from regions 2 and 3, as well as pure nickeloxide. Prior to the GITT measurement, all films were stabilized byundergoing 70 cycles in 1 M KOH electrolyte in the three-electrodecell. All measurements were carried out with a 20-s-long currentstep, followed by a relaxation time of 3580 s. The steady-state volt-age �SSV� was recorded when changes in the relaxation voltagewere less than 0.08 mV during 10 s.

The chemical diffusion coefficient was calculated from the volt-age response to a step current and the relaxation voltage response inan open circuit, using a relation based on the semi-infinite diffusionapproximation according to22

D̃ =4

� d2 ��Vs

ts �Vt

��t�2

ts �d2

D̃�14�

Here �Vs is the change of the steady-state equilibrium potential, Vtis the voltage response during the current step, ts is the step time,and t is the time. The relationship between voltage and time for thesemi-infinite approximation used in this model is

V�t���t 0 � t � ts �15�All voltage responses, obtained in this work upon application of acurrent step, show a linear dependence on the square root of the timeas predicted by Eq. 15. When we plotted the curve of the voltageresponse as a function of the square root of the time, the square ofthe correlation coefficient for a linear fit was of the order of 0.9995in all cases.

Effect of the experimental conditions.— The aim of this sectionis to show how sensitive the experimental setup is to the deinterca-lation process, even if the measurement is carried out in thediffusion-controlled regime. First, we show the response of the

Figure 12. Energy loss function obtained from measurements with p polar-ized light at 60° angle of incidence in the wavenumber region between 860and 360 cm−1. Three different voltages were applied corresponding to thefully bleached �−0.40 V�, intermediate �0.20 V�, and colored state �0.45 V�.Vibrations are assigned to �Ni–O�� and �O–H�l modes.

deintercalation of hydrogen under different current steps. We plotthe diffusion coefficient vs SSV as well as the dependence of theSSV on the extracted H+/Ni ratio. Figure 13 illustrates the behaviorof three polycrystalline films subjected to three different step cur-rents. We started the analysis when the films were in their bleachedstate in order to see changes in the proton deintercalation process. InFig. 13a-1, for the film sputtered in Ar + O2, it is possible to see one

minimum in D̃. In contrast to this, two minima in D̃ were observedfor the other two films in Fig. 13b-1 and c-1, deposited in Ar+ O2 + H2. For the first two films �Fig. 13a-1 and b-1�, the order ofmagnitude of the minima are around 10−17 cm2/s. This is differentfrom the result for the sample deposited with a high amount of O2

and H2 during sputtering �Fig. 13c-1�, which yielded D̃ of the orderof 10−14 cm2/s.

The appearance of a third minimum of D̃, around 10−13 cm2/s,can be discerned at a coloration potential of approximately 0.30 Vfor the two first two films �Fig. 13a-1 and 13b-1�. Because the po-tential saturates at this level, it is not possible to conclude whether itis a real minimum or an artifact of the measurements.

The differences between the curves of the SSVs at various cur-rent steps provide a criterion to evaluate the accuracy of the data. At

Figure 13. Chemical diffusion coefficient of H+ vs SSV �panels a1, b1, andc1�, and SSV vs atomic fraction of hydrogen �with respect to nickel� deinter-calated from the film �panels a2, b2, and c2�. Results are presented for threeNi1−xVxOy films deposited with the shown fractions of oxygen and hydrogenin the sputter plasma. Three current steps were applied: 0.1 mA �full circles�,0.3 mA �open circles�, and 0.6 mA �triangles�. In all measurements, a currentstep was applied during 20 s, followed by a relaxation time of 3580 s. Weused a 1 M KOH solution in a three-electrode configuration with a Ptcounter electrode and an Ag/AgCl reference electrode.

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high current steps, a relaxation time of 3580 s is not sufficient. Thefilm relaxes very slowly, which introduces an experimental limita-tion. We can take as accurate values the ones with overlapping SSVcurves for the lowest current steps, except for the part of the curve inFig. 13c-2 in which the change of the sign of the slope is probablyproduced by side effects.

Comparison between nickel oxide and nickel vanadiumoxide.— In nickel vanadium oxide, the vanadium is incorporated asan additive to facilitate film manufacturing by sputtering. It wasfound experimentally that V does not play any major role for theelectrochromism, i.e., the electrochemical and optical responseswere similar to those for pure nickel oxide. To further elaborate on

the possible influence of the vanadium, we present data on D̃ for apure hydrated nickel oxide film, made from a magnetic target of99.9% nickel in an Ar + O2 + H2 atmosphere �Fig. 14�, and a nickelvanadium oxide film deposited under similar conditions �Fig. 15�.

In order to provide reliable data on the diffusion coefficient, wefollowed the procedure described before in a more complex multi-step current experiment. The sequential current steps were 0.08,0.16, 0.24, 0.32, and 0.40 mA, in cycles of nine; the currents wereapplied for 20 s, and the voltages were recorded during a relaxationtime of 3580 s. The SSV was taken to be the one when the changewas less than 0.08 mV during 10 s.

The similarities between the data reported in Fig. 14 and 15 show

Figure 14. Chemical diffusion coefficient of H+ vs SSV �a� and SSV vsatomic fraction of hydrogen �with respect to nickel� �b� deintercalated from anickel oxide film deposited with the shown fractions of oxygen and hydrogenin the sputter plasma. Current steps of 0.08, 0.16, 0.24, 0.32, and 0.40 mA�nine steps each� were applied during 20 s, followed by a relaxation time of3580 s. We utilized a 1 M KOH solution in a three-electrode configurationwith a Pt counter electrode and an Ag/AgCl reference electrode.

that, effectively, the origin of the minima for D̃ lies solely in thebase material, i.e., in the hydrous nickel oxide. As seen in the po-tential curves for the two main changes of phase, we found well-defined plateaus at −0.1 and 0.23 V in both cases �cf. Fig. 14b and15b�.

In order to test the accuracy of the GITT method, we compared

with the information on D̃ calculated from the slope of the Randles-Sevčik plot �Fig. 4�, using the anodic peak, for the same film as theone reported on in Fig. 12a. The diffusion coefficient was 2.5

10−16 cm2/s, i.e., of the same order of magnitude as for D̃ deter-mined from the second minimum in Fig. 15a. The anodic peak atwhich the diffusion coefficient was determined and the SSV of thesecond minimum are at about the same electrochemical potential.

Discussion

The appearance of a sharp minimum in the diffusion coefficientvs potential can be understood as the result of a phase transitionduring the deintercalation of protons from the hydrated nickel-oxide-based films. The amplitude in the minima of the diffusioncoefficient is associated with the crystallinity, and well-crystallizedfilms exhibit deep minima upon extraction and insertion of charge.We can identify phases involved in the deintercalation process by

comparing the potentials at which the minima of D̃ appear and the

Figure 15. Chemical diffusion coefficient of H+ vs SSV �a� and SSV vsatomic fraction of hydrogen �with respect to nickel� �b� deintercalated from anickel vanadium oxide film deposited with the shown fractions of oxygenand hydrogen in the sputter plasma. Currents steps of 0.08, 0.16, 0.24, 0.32,and 0.40 mA �nine steps each� were applied during 20 s, followed by arelaxation time of 3580 s. We utilized a 1 M KOH solution in a three-electrode configuration with a Pt counter electrode and an Ag/AgCl refer-ence electrode.

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optical density �Fig. 6� because it is well known that the nickeloxy-hydroxide phases have higher optical absorption than the nickelhydroxide phases.1

It is seen that D̃ approaches zero as a discontinuity at the phasetransitions underlying the data in Fig. 14 and 15. This phenomenoncan be understood from the lattice gas model with interaction.48 Thechemical diffusion coefficient and the chemical potential �H+ for theFrumkin-type isotherm in this model are given by48,49

D̃ = U�1 − X�X� ��

�X� �16�

�H+ = �H+� + RgT0 ln� X

1 − X� + RgT0gX �17�

where X is the fraction of ions deintercalated due to the step current,U is the mobility at X = 0, �H+

� is the standard value of the chemicalpotential, and g is the interaction energy �divided by kBT0, where kBis Boltzmann’s constant� between the sites and includes the coulom-bic interaction and the strain field in the lattice. g is negative forattractive interaction and positive for repulsive interaction. Combin-ing Eq. 14 and 15, it can be shown that

D̃ = D̃X=0 + gD̃X=0�1 − X�X �18�

We do not have accurate measurements of the quantity of hydrogenpresent in the beginning, and hence we can only make a qualitative

analysis of the data presented in Fig. 13-15. At the two minima, D̃decreases by four to five orders of magnitude at the actual transition.This means that the model of the Frumkin-type isotherm, in whichthe electron insertion/extraction is slow enough that quasi-equilibrium conditions may be assumed, is an excellent approach for

a qualitative analysis.49 When D̃ approaches zero at the phasechange, the “critical” value of g is close to −4, which means that thenet force between the ions is attractive.50 A detailed review on thismatter can be found in Ref. 38.

Before discussing the results presented above, it is necessary tocomment on a number of factors that influence the measurementsand hence the extraction of the diffusion coefficients from the slopeof the voltage vs time. Even if the data were taken in the short-timedomain of finite space diffusion, there are some factors that wereneglected in the calculation of the diffusion coefficient, namely, theohmic drop, double-layer charging and kinetic effects due to thecurrent step, grain size distributions, etc. At least the three first ef-fects were minimized by choosing an adequate setup during themeasurements. Currently, there is much effort to include the men-tioned factors in order to produce better data on diffusion coeffi-cients, and different approaches have been developed to accomplishthat in measurements using potentiostatic intermittent titration andelectrochemical impedance spectroscopy.50,51 Corrected data, whichtake slow interfacial kinetics into account, have been published re-cently for lithium insertion in graphite.52 The corrected values re-ported in this latter work did not change the trends of the data inwhich different phases appeared. With regard to our own results, weexpect that changes in the magnitude of the reported data may ap-pear following refinements of the analysis, but the trends are ex-pected to remain unchanged.

Returning to the data in Fig. 14a and 15a, it is observed that thetwo first minima have their corresponding plateaus when the SSV isplotted vs the fraction of hydrogen intercalated into the films �seeFig. 14b and 15b�. This is the first indication that the observedminima are not an artifact due to the interfacial charge transfer ki-netics and/or an uncorrected contribution associated with ohmicdrops. The relaxation times were chosen to be sufficiently long toensure that the SSVs were close to the equilibrium, and the magni-tude and time of the step current were optimized in order that theexperiment was performed in the short time finite space diffusion

domain regime. Further, the magnitude of the chemical diffusioncoefficient was close to the one determined with the Randles-Sevčikmethod.

The minima for the diffusion coefficients as a function of thehydrogen content were identified and verified by comparison withoptical density data in the visible vs inserted charge �Fig. 6� and bycomparison with infrared absorption spectroscopy data for differentpotentials �Fig. 10 and 12�.

It is necessary to clarify some aspects concerning the changes inthe infrared absorption spectra at different potentials. Thus, in orderto understand the changes in the libration band, we need to explainthe changes in the stretching band �O–H��. The shift to lower fre-quencies of the stretching band observed in Fig. 10 is the result ofthe combination of the local external fields at the hydroxyl group ionsite created by the surrounding metal ions. This causes a polarizationof the O–H bond and decreases the strength of the bond.44,53 Thepolarization is enhanced by applying a positive potential, and thestrength of the bond is further decreased.45

The increase of the partial covalency of the Ni–O bond due to theelongation of the O–H bond increases the hybridization of the oxy-gen and nickel ions, thereby decreasing the bond length of the Ni–O.The libration band �O–H�l shifts to lower frequencies due to thecombination of the elongation of the O–H and the decrease of theNi–O bond. The spectrum measured at −0.40 V can be assigned tothe � phase of nickel hydroxide and the one at +0.20 V to the �phase because of the lower density of the latter phase. This confirmsthe interpretation of the first minimum in Fig. 14 and 15. It is wellknown that the � phase occurs at low contents of water and the �phase at high contents.14 The Ni�OH�2 structure of the � type con-sists of a random stacking of Ni�OH�2 layers with large interlayerseparation �0.76 nm�, whereas �-Ni�OH�2 has the same structureexcept that the interlayer separation is much smaller �0.46 nm�and hence the layers are well aligned.14

When the same film was measured at +0.40 V, the �O–H�l bandis further shifted to lower frequencies and a strong �Ni–O�� bandappears as a consequence of the extraction of a proton of the �phase of nickel hydroxide, thus being transformed to the � phase ofnickel oxy-hydroxide.45

The third minimum, if not an artifact, could be assigned to atransition from the �-nickel hydroxide to �-nickel oxyhydroxide,but no clear evidence for this could be found. The formulation ofthis mechanism only involves proton exchange, which in our case isthe main process governing the coloration. This assignment is fullyconsistent with the Bode reaction scheme given in Reaction 1, andwe note that different phases of nickel hydroxide and nickel oxyhy-droxide have been reported to exist in the following approximateranges of potentials involved in the dehydrogenation of pure nickelhydroxide and nickel hydroxide containing cobalt hydroxide to theirrespective phases of nickel oxyhydroxides:54 �-Ni�OH�2 for u� 0.15 V; �-Ni�OH�2 for 0.15 � u � 0.27 V; �-�NiO�OH for u� 0.30 V; and �-�NiO�OH for u � 0.34 V, where u is the potentialvs Ag/AgCl.

Recently, our results on ex situ grazing-incidence XRD, obtainedby use of synchrotron radiation, confirm the presence of hydroxideand oxy-hydroxide phases after stabilization in hydrated nickel-based oxides.55

Additional evidence for the phases present during the electro-chromic coloration was found from XPS analysis, as reportedbefore34 where it was shown that the intercalation takes place at thegrain surfaces so that the bulk remains NiO. In this earlier work,34

the changes in the Ni 2p core levels confirmed the transition Ni+2

→ Ni+3. The O 1s core levels showed the presence of the nickelhydroxide in the bleached state and the nickel oxyhydroxide in thecolored state. Changes in the O 1s core levels also pointed to modi-fications in the oxygen bonded with nickel. Because the coloration isdue to the extraction of protons, a mechanism analogous to the onefor the oxidation of nickel monoxide to dinickel trioxide through thecreation of a nickel +2 vacancy was proposed.34 The dehydrogena-

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tion of the nickel oxyhydroxide leaves a vacancy that is compen-sated by the creation of a hole in the nickel monoxide. This lasteffect could be related to the third minimum in Fig. 12 upon furtherdehydrogenation at higher voltages.

Conclusions

Nickel vanadium oxide thin films with good electrochromicproperties were produced by sputter deposition. The vanadium wasadded in order to quench the magnetism of pure nickel, therebyfacilitating deposition by magnetron sputtering. Optimized films,with high charge capacity and optical contrast, were deposited withintermediate concentrations of oxygen and hydrogen in the sputter-ing gas. Electrochemical and optical properties of the films werestudied by cyclic voltammetry, GITT, and optical spectroscopy.

Data on CV and CE, recorded on Pd-coated films, were similar tothose of uncoated films. Because only protons can move across thePd film, this shows that the coloration process is due to protontransfer. CVs showed that the maximum H/Ni ratio that can betransferred during intercalation was around 0.6. The proton diffusioncoefficient was measured as a function of the coloration potential byGITT. These data exhibit at least two distinct minima, which areconsistent with two phase changes within the intercalation process.These transitions can be interpreted, following the Bode reaction, asthe change from �-nickel hydroxide to �-nickel hydroxide and, sec-ond, from the �-nickel hydroxide to �-nickel oxyhydroxide. Diffu-sion coefficients obtained from the Randles-Sevčik equation were inagreement with the value at the second of these minima. The SSV ofthe second minimum and the potential of the anodic peak are atapproximately the same electrochemical potential. All of the phasetransitions were confirmed by infrared reflection-absorption spec-troscopy.

The optical absorption in the nickel vanadium oxide films re-mained small until the coloration potential corresponding to the sec-ond minimum was reached, at which point it increased rapidly. Thinfilms of hydrous nickel vanadium oxide and hydrous nickel oxidewere equivalent in their coloration mechanism, and the vanadium inthe films only increased the absorption in the bleached state but didnot play any role in the coloration. Thus, the electrochromism wasdue to the base material, i.e., nickel oxide.

Acknowledgment

The Swedish Foundation for Strategic Environmental Researchand the National Energy Administration of Sweden, through theÅngström Solar Center, partially supported this work.

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