journal of materials chemistry c6060 | j. mater. chem. c, 2019, 7 , 6059--6069 this journal is...

This journal is©The Royal Society of Chemistry 2019 J. Mater. Chem. C, 2019, 7, 6059--6069 | 6059

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7, 6059

Cationic compositional effects on the bias-stressstabilities of thin film transistors using In–Ga–Zn–Ochannels prepared by atomic layer deposition

Seung-Bo Ko, a Nak-Jin Seong,b Kyujeong Choi,b So-Jung Yoon,a Se-Na Choia

and Sung-Min Yoon *a

Cationic compositional effects of amorphous In–Ga–Zn–O (a-IGZO) prepared by atomic layer

deposition (ALD) were strategically investigated for thin film transistor applications. The atomic

compositions (In : Ga : Zn) of ALD-IGZO films were varied to 1 : 1 : 1, 1 : 1 : 3, and 1 : 1 : 5 by controlling the

ALD cyclic ratios. The relative content of oxygen vacancies and temperature-dependent electrical

conductivities among the films markedly varied with the In/Zn ratio. The device employing the 1 : 1 : 5

composition exhibited inferior characteristics owing to the excessive Zn content in the IGZO channel.

With increasing In/Zn ratio, the density of subgap states near both the conduction and valence bands

increased, resulting in a higher degree of bias-stress instability. The device employing the 1 : 1 : 3

composition exhibited the most promising device characteristics including excellent stabilities under

positive bias-stress at 60 1C and a negative bias-illumination-stress condition using a green wavelength,

in which the threshold voltage shifts were estimated to be as low as +0.8 and �1.5 V, respectively.

1. Introduction

Recently, amorphous In–Ga–Zn–O (a-IGZO) materials havebeen mainly employed as a channel material for the backplaneTFTs of flat panel displays (FPDs) owing to their superiorcharacteristics of excellent uniformity, high on/off currentratio, and superior carrier mobility to other amorphous oxidesemiconductors (AOSs).1,2 Moreover, a-IGZO thin films havebeen actively researched for high-end future electronic applica-tions such as flexible displays,3,4 power devices,5 sensors,6,7

thermoelectric devices,8,9 and logic architectures, includingmemory devices.10,11 Conventionally, IGZO films have beenprepared by a sputtering deposition technique.12 However,the sputtering process may not be appropriate for the for-mation of IGZO thin films, especially for such applicationsdemanding a lower process temperature, excellent film con-formality, and precise control of film thickness and composi-tion, because the sputter-deposited IGZO films typically requirea post-annealing process at higher temperature to reduceplasma-induced damage. Accurate thickness and compositioncontrol of the IGZO thin films is also very difficult due to thecontinuous consumption of the sputtering target with time

evolution, especially for the cases of using a single target ofmulticomponent compositions. On the other hand, the atomiclayer deposition (ALD) method is based on plasma-free gas-phase chemical reactions, resulting in better film quality evenat a lower deposition temperature. Furthermore, owing tothe self-limiting growth mechanism, the film thickness andcomposition can always be precisely controlled at the atomicscale with excellent conformality and higher film density.13,14

Alternatively, a long process time, considered as a criticalweakness of ALD, is not such a main concern, because thethickness of the active channel layer used in TFTs is at most afew tens of nanometers. Therefore, by exploiting these uniqueadvantages of ALD, high quality a-IGZO thin films can beapplied to future electronics demanding complex device geo-metry or mechanical flexibility.

Nevertheless, there are few reports on IGZO thin filmsprepared by ALD because of process complexities and relatedtechnical considerations for forming high-quality quaternaryoxide films. The growth of constituent layers in a-IGZO thinfilms may be hindered by non-ideal surface thermodynamicsand kinetics between metal precursors and oxidants.15 More-over, the ALD process window becomes narrow owing to thedifference in vaporization temperatures among metal precursors.14

To overcome these barriers, two experimental strategies wereestablished. One is to employ a two-metal precursor system andthe other is to develop a new bimetallic In–Ga single precursor.In other words, a two-metal precursor system composed of an

a Department of Advanced Materials Engineering for Information and Electronics,

Kyung Hee University, Yongin, Gyeonggi-do 17104, Korea.

E-mail: [email protected] NCD Co., Ltd, Daejeon 34015, Korea

Received 1st March 2019,Accepted 15th April 2019

DOI: 10.1039/c9tc01164a

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Journal ofMaterials Chemistry C

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6060 | J. Mater. Chem. C, 2019, 7, 6059--6069 This journal is©The Royal Society of Chemistry 2019

In–Ga single precursor (In–Ga) with an atomic ratio of 1 : 1 anddiethyl zinc (DEZn). We have previously reported IGZO deposi-tion temperature dependent device characteristics, including thebias stability of ALD-grown IGZO channel TFTs,16 especiallywhen the channel thickness was reduced to 6 nm.17 However,the ultimate goal is to enhance both the characteristics of deviceperformance and bias stability of ALD-IGZO channel TFTs.

One powerful solution is to control the cationic compositionsof the ALD-IGZO channel layers. The device characteristics andbias stability of IGZO TFTs are closely dependent on the subgapdensity of states (DOS) divided into shallow donor, localizedtrap, and fully-occupied deep states.18–20 The subgap DOS can beexpected to be effectively modulated by changes in the cationiccompositions. Thus, in this work, ALD sub-cycles were designedand controlled for the formation of IGZO channels so that theinfluence of the channel compositions, especially the In/Zn ratio,on the electronic nature and device characteristics includingbias stability could be elucidated. Even though there havebeen many previous studies on the compositional effects ofconventional sputter-deposited IGZO channels on the devicecharacteristics,21–23 the ALD-IGZO channels prepared by layer-by-layer deposition techniques may show totally differentbehaviors because of modulations in the bonding nature anddefect states compared with those of the conventional sputter-deposited films. Furthermore, investigations on the effects ofchannel compositions related to the ALD cyclic ratios on thedevice reliability, which is one of the most important specifica-tions to employ ALD-IGZO TFTs in the aforementioned novelapplications, have never been reported so far.15,24,25 In fact, wediscovered very interesting phenomena attributed to the ALDprocess and promising TFT characteristics showing compar-able device performance with robust stability to those ofthe sputter-deposited IGZO TFTs. Therefore, this pioneeringresearch will emphasize the significance of compositionalcontrol of the IGZO channel and extend the ALD process forIGZO into cutting-edge future applications.

2. Experimental

First of all, the IGZO thin films were prepared by the ALDmethod to elucidate the composition dependent physical char-acteristics. An In–Ga single precursor, DEZn, and ozone (O3)were employed as In–Ga, Zn source, and oxidant, respectively.The ALD process temperature was set at 150 1C, which had beenverified as the most optimum condition.16 The ALD cyclic ratios(In–Ga : Zn), a key process parameter for our ALD process, weremodulated to be 4 : 1, 2 : 2, and 2 : 3, as shown in Fig. 1(a);for convenience, each thin film was termed A, B, and C,respectively. One-hundred ALD super-cycles were applied toform a 20 nm-thick ALD IGZO layer.

Fig. 1(b) shows a schematic cross-sectional diagram of thetop-gate-bottom-contact (TGBC) IGZO TFTs using the ALD-grownIGZO channel layers with various compositions. First, thesource/drain (S/D) electrodes were patterned on 150 nm-thickindium tin oxide (ITO)-coated glass substrates by wet chemical

etching. The IGZO active channel layers were deposited by ALDat 150 1C, in which the cationic compositions of the IGZOchannel layers were varied as A, B, and C. For convenience, eachdevice was termed Dev. A, Dev. B, and Dev. C. Then, 9 nm-thickAl2O3 thin films were prepared as protection layers (PLs) by ALDand were simultaneously patterned with the IGZO active layersby using a diluted hydrofluoric acid-based (DHF) wet etchant.Then, 100 nm-thick Al2O3 films were deposited as gate insulators(GIs) by ALD and contact holes were formed to electricallyconnect the S/D electrodes by using a phosphoric acid-basedwet chemical etchant. The deposition processes for the Al2O3 PLsand GIs were performed at 150 1C by using trimethylaluminum(TMA) and H2O as the aluminum precursor and oxygen reactant,respectively, in which the deposition rate was estimated to beapproximately 12.5 Å per cycle. Finally, 150 nm-thick ITO thinfilms were deposited by DC-magnetron sputtering at roomtemperature (RT) and patterned by using a lift-off process toform gate electrodes and S/D pads. All the fabricated deviceswere post-annealed in a rapid thermal annealing system at180 1C for 1 h in an oxygen environment. The channel width/length (W/L) of the examined devices was 40/20 mm.

Thickness values of the ALD-grown IGZO films were mea-sured by high-resolution field-emission scanning electronmicroscopy (FE-SEM). The oxygen vacancies and cationiccompositions in the ALD-grown IGZO films were examined byX-ray photoelectron spectroscopy (XPS) analysis with depth-profiling. The temperature-dependent electrical conductivity(sc) of the IGZO thin films with various compositions was alsoinvestigated to elucidate the cationic compositional effects onthe device characteristics. The bias, illumination, temperature

Fig. 1 (a) Schematic illustrations of the designed ALD sub-cycles forvaried ALD-IGZO cationic compositions of films A, B, and C. 100 super-cycles were employed to form 20 nm-thick films. (b) Schematic cross-sectional view of the fabricated ALD-IGZO TFT.

Paper Journal of Materials Chemistry C


stress instabilities and temperature-dependent transfer char-acteristics of the fabricated devices were measured by using asemiconductor parameter analyzer (Keithley 4200A-SCS) in adark box or chamber-type probe station equipped with hotchunk. The device parameters including carrier mobility ofeach device were accurately extracted by using the relativedielectric constant of 9.7 for 100 nm-thick ALD-Al2O3 thin films,which was calculated from the capacitance–frequency charac-teristics obtained for the metal–insulator–metal capacitors at afrequency of 1 MHz.

3. Results and discussion

The thickness values of all the ALD-IGZO films prepared withdifferent ALD cyclic ratios were estimated to be approximately17–18 nm, which was slightly thinner than the value estimatedfrom the growth per cycle of ALD process. However, theobtained thickness values were confirmed to be sufficientlyuniform for considering the film thickness variations amongthe ALD-IGZO active layers with different cationic composi-tions. Prior to investigations on the device characteristics, thephysical properties of the ALD-IGZO films with various atomiccompositions were examined. Fig. 2(a)–(c) show the XPS depth-profile analyses for estimating the cationic compositions of theIGZO films prepared with cyclic ratios of 4 : 1, 2 : 2, and 2 : 3,respectively. The film compositions remain constant in thethickness range of all the deposited films. The atomic ratios(In : Ga : Zn : O) of films A, B, and C were found to be approxi-mately 1 : 1 : 1 : 3.5, 1 : 1 : 3 : 5, and 1 : 1 : 5 : 7.5, respectively.Table 1 summarizes the calculated atomic compositions andIn/Zn ratios for the prepared IGZO films. It was found fromthese results that the contents of both In and Ga equallyincreased with increasing the relative cycle number of theIn–Ga precursor, as expected, and that the In/Zn ratio alsoincreased from 0.2 to 1.0. Thus, the choice of the In–Ga singleprecursor was suggested as appropriate for controlling thecationic compositions of the ALD IGZO films.

Next, to investigate the bonding nature of the oxygen-relateddefects in the IGZO films prepared with cyclic ratios of 4 : 1,2 : 2, and 2 : 3, which was expected to be highly related to the

electronic nature of the IGZO films, oxygen 1s (O 1s) spectra ofeach film were measured using XPS, as shown in Fig. 3(a), (b),and (c), respectively. The O 1s spectra can be deconvoluted intothree sub-peaks, metal–oxygen bonding (peak LE), oxygenvacancies (VO, peak ME), and contamination on the surface(peak HE), respectively, based on the binding energy withoxygen. Deconvolution of the O 1s peak spectra obtained at adepth after 30 s etching from the surface, at which the cationiccompositional ratios were defined, was conducted to preciselyfind out the accurate relative peak areas in the specifiedcompositions and to minimize the surface contamination.Thus, we can neglect the surface contamination term, peakHE. Additionally, the deconvolution process was carefullycarried out to enhance the accuracy of analysis. The peaksshowed a Shirley-type background and the background bindingenergy was set from 527 to 534 eV to generate accuratestandards. The peaks LE and ME were decomposed to be530.1 and 531.3 eV, respectively. The full width at half-maximum parameter and Lorentzian–Gaussian percentagewere set as 1% and 1% for the conventional s orbital, respec-tively. The calculated relative area fractions of the deconvolutedpeaks LE and ME for the three films with respect to total areaare summarized in Table 2. Among them, the peak ME corres-ponding to VO was focused on from the viewpoint that varia-tions in the amount of VO play the biggest role in determiningthe device characteristics.26,27 The relative area ratio of peakME increased in the order of films B, A, and C. In other words,film C contained a higher amount of VO and, hence, included arelatively larger number of conduction carriers and/or trapstates owing to a higher content of Zn among the films. Therewere also meaningful variations between the films A and B.Although film A contained a larger content of incorporated Ga,

Fig. 2 Estimated cationic compositions of the ALD-IGZO films (a) A, (b) B, and (c) C as a function of etch time by XPS depth profile analysis.

Table 1 Summary of the atomic compositions of three IGZO thin filmsprepared on SiO2/Si substrates. Atomic percent of each ALD-IGZO filmwas calculated at a point after an etch time of 30 s in the XPS depth profiles

Film ALD sub-cycle Atomic percent (In : Ga : Zn : O) In/Zn ratio

A In–Ga : DEZn = 4 : 1 15 : 18.5 : 14 : 52 : 5 (1 : 1 : 1 : 3.5) 1B In–Ga : DEZn = 2 : 2 8.5 : 9.5 : 30.5 : 51.5 (1 : 1 : 3 : 5) 0.3C In–Ga : DEZn = 2 : 3 6.7 : 6.6 : 35.9 : 50.8 (1 : 1 : 5 : 7.5) 0.2

Journal of Materials Chemistry C Paper


acting as a suppressor for the formation of oxygen vacancies,than film B, a higher content of oxygen vacancy in film A couldbe explained by an optimum In content and In/Zn ratio for theformation of In cation defects or In–Zn defects in film A, becausethese defects can be main origins of oxygen vacancies.19,20 Thus,the variations in oxygen vacancy content for both films weresuggested to be determined by the two competing factors ofoxygen vacancy suppressors and generators. Consequently,the cationic compositions were found to have notableimpacts on the bonding natures and their variations instructural defects including the VO within the ALD-IGZOthin films. The physical relationship between the bondingnatures and the device characteristics including the biasstability of the ALD-IGZO TFTs will be discussed below in adetailed way.

Fig. 4 shows the variation in electrical conductivity (sc) ofthe ALD-IGZO films with three compositions as a function ofmeasurement temperature from 50 to 200 1C in the forward andreverse directions, in which the specimens were prepared oninsulating SiO2/Si substrates. The ramping rate and measure-ment speed were set as 5 1C min�1 and 1 pts per s, respectively.Considering that it is hard to directly perform Hall measure-ments owing to difficulties in forming ohmic contacts for themeasurements and the relatively low conduction carrierconcentration of the IGZO compositions, this measurementcan be expected to provide us with not only detailed informationabout the intrinsic electronic natures of the ALD-IGZO thin filmsbut also unique information on the temperature-dependentcarrier transport and donor-state formation processes18 whenthe cationic compositions were varied for the IGZO films pre-pared by ALD. The values of sc for the IGZO films well obey theArrhenius relationship described in eqn (1), where s0 is a

prefactor, Ea is activation energy, k is the Boltzmann constant,and T is absolute temperature.

sc = s0�exp(�Ea/kT) (1)

Thus, the slopes of the sc–1/T curves shown in Fig. 4 corre-spond to the Ea for carrier transport of electrical conduction.The temperature dependent sc variations could be divided intothree temperature regions, I, II, and III, according to thetransitions of the sc variations with temperature evolution.For region I, the values of sc for all the films monotonouslyincreased with increasing temperature, which clearly suggestedthat the electrical conduction of the IGZO films is a thermally-activated process, showing that the ALD-IGZO films havetypical semiconducting properties. Alternatively, for region II,the changes in Ea were different for the three films. In otherwords, while the Ea for films A and B decreased compared withthose estimated in region I, for film C, the Ea showed anincreasing trend in region II compared to that in region I.These variations in Ea were suggested to be closely associatedwith the formation of oxygen-related defects. The decrease in Ea

for films A and B in region II can be attributed to the higherIn/Zn ratios than that of film C. With increasing In contents

Fig. 3 Variations in the oxygen 1s (O 1s) XPS spectra obtained after 30 s surface etching with deconvoluted peaks for the ALD-IGZO films (a) A, (b) B, and(c) C employing different cationic compositions by controlling ALD cyclic ratios.

Table 2 Comparison of the calculated relative area fractions of thedeconvoluted peaks, LE and ME, for the three IGZO films with respect tototal area. Peaks LE and ME correspond to metal–oxygen bonding andoxygen vacancies (VO) in the ALD-IGZO films, respectively

B.E. [eV] Film A Film B Film C

O 1s LE 530.1 79.0 82.4 66.8O 1s ME (VO) 531.3 21.0 17.6 33.2

Fig. 4 Variations in electrical conductivity (sc) of the ALD-IGZO films withthree compositions as a function of measurement temperature using anArrhenius plot. The temperature was swept from 50 to 200 1C in theforward and reverse directions. The ramping rate and measurement speedwere set as 5 1C min�1 and 1 pts per s, respectively. The dashed linescorrespond to the boundaries of temperature regions I, II, and III showingthe transitions of sc variation with temperature evolution.



within IGZO, the oxygen vacancies formed by edge-sharing InO6

octahedrons, which act as carrier generators,28 increase andgenerate more thermally-activated donor electrons. On theother hand, the increase in Ea of film C in region II can besuggested to be caused by higher Zn contents. The oxygenvacancies acting as trap sites for the thermally-activated freeelectrons may easily form with the increase in the number ofcorner-sharing ZnO4 tetrahedrons.28 As a result, carrier trans-port is retarded by the defect scattering related to the electrontrapping. Local formation of ZnO nano-crystallites might beanother origin for the increase in Ea for film C.23 Furthermore,the formation energies of trap states and the degree of localcrystallization are expected to decrease and increase, respec-tively, in region II. It was found from these results that thetemperature-dependent carrier transport behaviors were alsoclosely related to the variations in In/Zn ratios among the threeALD-IGZO thin films prepared by using different sub-cycleratios. Finally, for region III, the sc values of all the films didnot change even when the measurement temperature inverselydecreased, which was determined by thermal treatment at200 1C. It was interesting to note that the finally obtained sc

showed higher values in the order of films A, B, and C. This canbe attributed to the many InO6 octahedral edge-sharing siteswith increasing In/Zn ratio. Consequently, the cationic compo-sitions of the ALD-IGZO films were found to have criticalimpacts on the relative amounts and distributions of oxygen-related defects, resulting in meaningful modulations in theelectronic natures of the ALD-IGZO thin films.

Based on the physical properties of the ALD-IGZO filmswith different cationic compositions, as mentioned above, thedevice characteristics of the IGZO TFTs were investigated.

Fig. 5(a)–(c) show the drain current (IDS)–gate voltage (VGS)transfer characteristics and gate leakage current (IGS) ofDev. A, B, and C, respectively. The VGS was swept from �20(especially �30 V for Dev. C) to 20 V in the forward and reversedirections and the drain bias (VDS) was set at 10.5 V. Dev. A andDev. B exhibited great device-to-device uniformity over 9 pat-terns and no hysteresis behavior in their transfer characteris-tics. On the other hand, the transfer curves of Dev. C showedterrible fluctuations with hysteretic behavior among the9 devices. Considering that the IGZO channel for Dev. Ccontained a higher content of oxygen vacancies (Fig. 3(c)), theinferior device-to-device uniformity might mainly result fromthe excessively-included oxygen vacancies and the large varia-tions in their distributions among devices. Fig. 5(d) comparesthe transfer characteristics among the devices. The deviceparameters including the threshold voltage (VTH), the carriermobility in the saturation region (msat), and the subthresholdslope (SS) value, which were extracted from the transfer curvesare summarized in Table 3. It is noteworthy that the fabricateddevices employing different ALD-IGZO cationic compositionsshowed remarkable variations in their transfer characteristics.Above all, Dev. C showed inferior characteristics compared tothose of Dev. A and Dev. B.

As expected from the O 1s analysis and temperature-dependent sc measurement, the negative turn-on characteris-tics and too-large SS values could originate from the higherdensity of carriers and trap sites within the IGZO activechannel. When the Zn content in the IGZO layer exceeds acritical concentration, grain boundaries can be generated dueto the local crystallization of ZnO with an increase in conduc-tion carriers and trap sites, and hence, msat and SS were found

Fig. 5 IDS–VGS transfer characteristics of the fabricated Dev. (a) A, (b) B, and (c) C, respectively. (d) Comparison of the transfer characteristics among thedevices employing different IGZO channel compositions. VGS was swept in the forward and reverse directions and the drain bias (VDS) was set at 10.5 V.



to be markedly reduced and deteriorated, respectively, owing tothe grain-boundary and defect scattering retarding the carriertransport. Alternatively, there were meaningful differences inboth VTH and msat, which might be due to the different In/Znratios, between Dev. A and Dev. B. However, the SS of the twodevices exhibited similar values. Therefore, to elucidate theseappreciable and subtle differences between the two devices,in-depth analyses on the device characteristics need to beperformed.

To figure out the detailed differences in device behaviorbetween Dev. A and Dev. B, the subgap DOS in the IGZOchannels was extracted. From the calculated subgap DOS, wecan estimate the relative distributions of shallow donor-states,tail-states, and localized trap-states within the bandgap, whichinfluence the transfer characteristics of the fabricated devicesemploying ALD-IGZO channels with different cationic composi-tions of 1 : 1 : 1 and 1 : 1 : 3. The subgap state extraction methodscan be mainly classified into three types: C-V measurement,29

temperature dependent field-effect method,30 and computersimulations.31 In this work, we adopted the temperature depen-dent field-effect method using the Meyer–Neldel (MN) rule.30,32

The temperature-dependent IDS of the ALD-IGZO TFTs wassupposed to vary with a thermally-activated process, as shownin Fig. 4. Thus, it can also be described by the Arrhenius relationin eqn (2), where IDS0 is a prefactor for IDS, Ea is activation energy,k is the Boltzmann constant, and T is absolute temperature.

IDS = IDS0�exp(�Ea/kT) (2)

Assuming that the ALD-grown amorphous IGZO thin filmscorrespond to a disordered system, which frequently appearsin a group of inorganic semiconductor materials,33 the prefactorIDS0 and Ea can be described by eqn (3) obeying the MN rule,where IDS00 is a prefactor for IDS0, and EMN is the MN energy.

IDS0 = IDS00�exp(Ea/EMN) (3)

Combining eqn (2) and (3), we can derive the IDS with thefollowing eqn (4).

IDS ¼ IDS00 � exp1

EMN� 1

kT

� �Ea

� �(4)

Fig. 6(a) and (b) show the Arrhenius plots derived from eqn (4)as a function of 1/kT at various values of VGS, in which theintersections in x- and y-coordinates represent the EMN andIDS00 values, respectively. The formation of definite inter-sections among the extrapolated lines even with different Ea

values means that the fabricated ALD-IGZO TFTs well obey the

MN rule. The subgap DOS of our evaluated ALD-IGZO TFTs canbe extracted from those parameters by using Poisson equationswith 0 K Fermi statistics.30 This method is a very simple andaccurate process to extract the subgap DOS because it onlyrequires temperature-dependent field-effect measurements andprovides verification procedures for the calculated results.Fig. 6(c) compares the extracted density of subgap states fortwo devices (A and B) from the conduction band minimum (EC)down to 0.7 eV. This range was supposed to be enough toexplain the differences in electrical properties of the twodevices because the trap states deeper than 0.8 eV from EC donot influence the carrier transport and the turn-on character-istics of the AOS TFTs.18 The extracted subgap DOS provides uswith two important observations as follows: (1) the DOS values

Table 3 Summary of extracted device parameters from the IDS–VGS

characteristics of the ALD-IGZO TFTs employing different ALD-IGZOchannel compositions

Device In : Ga : ZnIn/Znratio VTH [V]

msat

[cm2 V�1 s�1]SS[V dec�1]

A 1 : 1 : 1 1 �3.26 � 0.32 15.1 � 0.31 0.17 � 0.02B 1 : 1 : 3 0.3 0.06 � 0.18 11.3 � 0.24 0.14 � 0.02C 1 : 1 : 5 0.2 �8.04 0.04 5.98

Fig. 6 Arrhenius plots as a function of 1/kT at various VGS in the sub-threshold region and extracted parameters for Dev. (a) A and (b) B, whichwere obtained by measuring the IDS–VGS transfer curves of the ALD-IGZOTFTs at various temperatures (25, 35, 45, and 55 1C) at VDS = 0.5 V.(c) Comparison of the extracted subgap DOS near the Ec of Dev. A and Bemploying different channel compositions.



calculated at the EC for Dev. A and Dev. B were estimated to beapproximately 9.21� 1017 eV�1 cm�3 and 3.59� 1017 eV�1 cm�3,respectively. This meaningful difference could be a main originof the lower VTH and higher msat of Dev. A, since the VTH and msat

are supposed to be affected by donor-state densities located justbelow the EC; (2) the density of tail states of Dev. A was estimatedto be higher than that of Dev. B, as indicated by a dotted circle inFig. 6(c). The tail states located in shallow levels in the E–Ec

range of 0.1 to 0.0 eV,18 which are typically composed ofinterfacial trap and bulk trap densities, are regarded as closelyrelated to SS and positive bias-stress (PBS) instability of the IGZOTFTs. Considering that the SS is mainly dependent onthe defects located at the insulator/semiconductor interface,the total amount of interfacial trap density might be almost thesame for Dev. A and Dev. B. Thus, a higher density of tail statesfor Dev. A can be assumed to be due to more bulk trap states,which are induced by internal bonding originating from thecompositional variations. As a result, these variations causemarked differences in PBS instabilities between the devices.From these subgap DOS extraction procedures, we can defi-nitely conclude that a higher In/Zn ratio could result in a highermsat and lower VTH, and more bulk trap sites for Dev. A owing tothe higher densities of the donor and tail states near theconduction band. These results are in good agreement withthe theoretical investigation that found that In–Zn defects aremainly responsible for the subgap states in the upper half ofthe band gap.19 Consequently, the cationic compositions of theALD-IGZO channel layers have great impacts on the deviceperformance of the fabricated IGZO TFTs.

The next concern is the operation reliability under givenstress conditions, because acceptable device reliabilities ofALD-IGZO TFTs have hardly been reported and the feasible

origins for the instabilities have rarely been explored so far. Tounderstand the device operation stability of the fabricatedIGZO TFTs, PBS, positive bias-temperature-stress (PBTS), andnegative bias-illumination-stress (NBIS) conditions are gener-ally imposed during device operation. In this work, the PB(T)Sand NBIS instabilities of Dev. A and Dev. B were compared toinvestigate the device reliability of the ALD-IGZO TFTs and thechannel compositional effects on the origins of the operationinstabilities. Fig. 7(a), (b) and (c), (d) show the variations in thetransfer curves with a lapse of stress time for 104 s under thePBS and PBTS conditions at 60 1C for Dev. A and Dev. B,respectively. A VGS bias of +20 V was applied during the testsand the VDS was set at 10.5 V. The values of threshold voltageshift (DVTH) estimated after the stress time of 104 s aresummarized in Table 4. Commonly, PB(T)S instability is relatedto electron trapping at the insulator/semiconductor interfaceand/or within the bulk insulator.34 While the DVTH values ofDev. B were estimated to be +0.5 and +0.8 V under the PBS andPBTS conditions, respectively, for Dev. A, the DVTH values wereobtained as �0.4 and �4.1 V under the same conditions. Here,we have some noticeable points to be discussed as follows:(1) the PB(T)S stabilities for both devices are sufficiently com-parable or superior (especially for Dev. A) to those of the best

Fig. 7 Variations in transfer curves with the lapse of stress time for 104 s under PB(T)S conditions for Dev. (a) A and (b) B at RT and for Dev. (c) A and (d) Bat 60 1C. The VGS bias of +20 V was applied during the tests and the VDS was set at 10.5 V.

Table 4 Summary of threshold voltage shift (DVTH) values estimated aftera lapse of stress time for 104 s for two devices employing different ALD-IGZO channel compositions under PBS, PBTS, and NBIS conditions

Device In : Ga : ZnIn/Znratio

PBSDVTH [V]

PBTSDVTH [V]

NBIS greenDVTH [V]

NBIS blueDVTH [V]

A 1 : 1 : 1 1 �0.4 �4.1 �5.4 �15.1B 1 : 1 : 3 0.3 +0.5 +0.8 �1.5 �14.8



devices using sputter-deposited IGZO active channels; (2) forDev. B, the value of DVTH smaller than +0.8 V even at 60 1Ccould be well explained by the simple electron trapping model,considering that the msat and SS values did not experiencemarked changes during the tests; (3) for Dev. A, it is noteworthythat the VTH anomalously shifted in a negative direction underthe PBS and PBTS conditions, which could not be simplyexplained by any charge-trapping model. There have been afew previous studies reporting a negative DVTH under thePB(T)S condition for devices using sputter-deposited IGZOchannel layers.35–37 Y. H. Chang et al. reported that IGZO TFTsemploying ALD-Al2O3 GIs prepared at a lower temperature of120 1C showed a negative DVTH owing to the breakage ofresidual Al–OH bonds and subsequent migration of hydrogensinto the IGZO channel layer under the PBS condition.36

However, since Dev. B and the sputter-deposited IGZO TFT(not shown in here) employing the same Al2O3 GIs prepared at150 1C showed a positive DVTH under the PB(T)S condition,hydrogen migration could be excluded. The formation of ALDsub-cycles designed to accomplish the various film composi-tions of the ALD-IGZO thin films was suggested to be the mostfeasible scenario to explain the anomalous negative DVTH

of Dev. A. In other words, as presented in Fig. 1(a), to preparecomposition A, only one DEZn sub-cycle is inserted every 4

sub-cycles of the In–Ga precursor. Thus, the adsorption of Znand the formation of ZnO can be so localized that the overallamorphous network of IGZO cannot be suitably formed, eventhough the cationic composition is keep as 1 : 1 : 1. Fig. 8schematically illustrates the feasible mechanism for the freeelectron generation in Dev. A under the PB(T)S condition. Inother words, some weak and unstable metal–oxygen sitesremain within the deposited film (Fig. 8(a)). If the weak bondsare broken by thermal vibration due to the lowered bondingstrength with oxygen (Fig. 8(b)), the generated free electronsinduce the negative DVTH, which can be accelerated at a highertemperature owing to larger lattice vibration (Fig. 8(c)).As a result, for Dev. A, the free electron generation can bemore dominant than the conventional electron trapping byinducing the negative DVTH under the PB(T)S condition, asshown in Fig. 8(d). Alternatively, as shown in Fig. 8(e), electrontrapping is expected to dominantly occur in Dev. B underthe PB(T)S condition. However, this scenario should becarefully checked because these comparisons with onlytwo compositions cannot be sufficient to accurately investi-gate the impacts of film compositions and ALD cycleformations on the device stability. Therefore, not only experi-mental analysis including the systematic decomposition ofthe subgap DOS,38 but also theoretical approaches to defect

Fig. 8 (a)–(c) Schematic illustrations for the feasible mechanism of the bond breakage and free electron generation in Dev. A under PB(T)S conditions.The bonds indicated by a dotted line in (a) mean weak bonds. Black and grey polygonal lines in (b) and (c) represent the thermally-activated latticevibrations for normal and weak bonds, respectively. The schematic band diagrams for the PB(T)S instability mechanism of (d) Dev. A and (e) Dev. B.



structures19,26 for various other cyclic sequences will beperformed as future works.

Fig. 9(a), (b) and (c), (d) show the variations in the transfercurves with a lapse of stress time for 104 s under the NBIScondition employing green and blue light for Dev. A and Dev. B,respectively. The wavelengths of the green and blue light werefixed at 530 and 470 nm, respectively, and the power intensitywas set at 0.1 mW cm�2. A VGS bias of �20 V was applied duringthe tests and the VDS was set at 10.5 V. Typically, NBISinstability is related to fully occupied deep neutral oxygenvacancy states (V0

O) near the valence band maximum (Ev). Underthe NBIS condition, the V0

O states close to the Fermi Level (EF)can be photo-excited into VO

2+ and 2e� by photon energy, andhence, the VTH can be negatively shifted due to the increase inconduction electrons.26,39 As summarized in Table 4, the DVTH

of Dev. A and Dev. B under the NBIS condition using the greenwavelength was estimated to be �5.4 V and�1.5 V, respectively.It was also found that the DVTH under both illuminationconditions using green and blue wavelengths was larger forDev. A than Dev. B. The higher density of V0

O states near the Ev

for Dev. A using the IGZO channel with a higher In/Zn ratio,which can generate a larger number of VO

2+ and free electrons,contributed to the larger DVTH. In other words, the subgapDOS extraction and NBIS results suggest that the donorand trap states near the Ec, and V0

O states near the Ev increasewith increasing the In/Zn ratio within the IGZO compositions.Consequently, the cationic compositions of the ALD-IGZOactive channels have a significant impact on the bias stabilityas well as the device characteristics because of the markeddifferences in the DOS within the bandgap for the ALD-IGZOthin films.

These results sufficiently show the cationic composition-dependent device performance of the ALD-IGZO TFTs. How-ever, more systematic and theoretical investigations should beperformed with various other compositions to properly specifythe comprehensive channel composition effects including therole of Ga in the TFT characteristics employing the ALD-IGZOactive channel layers, because the physical and electricalproperties of IGZO thin films are sensitively dependent onthe bonding nature of oxygen with the cations in amorphousstructures according to channel compositions and ALD cyclicsequences. Furthermore, we compared our Dev. A and B withthe conventional sputter-deposited IGZO TFTs with TGBCstructures and ALD-Al2O3 GIs.39–41 Although TFT performanceis markedly influenced by detailed conditions of device andprocess designs, these comparisons apparently revealed thatour ALD-IGZO TFT showed better performance including devicestabilities comparable to those of the devices obtained usingthe sputter-deposited IGZO channel layers.

4. Conclusions

The effects of IGZO channel composition were elucidated toenhance the device performance including the robust biasstability of ALD-IGZO TFTs. The channel compositions (atomicratio of In : Ga : Zn) of ALD-IGZO were varied as 1 : 1 : 1, 1 : 1 : 3,and 1 : 1 : 5 by controlling the ALD cyclic ratios. The relativecontent of VO varied with the compositional variations, whichwas attributed to the compositional dependences of the bond-ing nature and structural defects within the ALD-IGZO thinfilms. The behavior of electrical conduction was found to vary

Fig. 9 Variations in transfer curves with the lapse of stress time for 104 s under NBIS conditions for Dev. (a) A and (b) B under illumination with a greenwavelength and for Dev. (c) A and (d) B under illumination with a blue wavelength. The wavelengths of green and blue light were fixed at 530 and 470 nm,respectively, and the power intensity was set at 0.1 mW cm�2. The VGS bias of �20 V was applied during the tests and the VDS was set at 10.5 V.



with In/Zn ratio. It was found that the physical properties of theALD-IGZO thin films were significantly affected by the cationiccompositions.

The composition-dependent electrical properties of theALD-IGZO thin films were reflected in the device characteristicsemploying various IGZO channel compositions. The deviceemploying the 1 : 1 : 5 composition showed inferior character-istics due to a too high VO density and defect scattering withinthe IGZO channel. Temperature-dependent field-effect mea-surements were performed to extract the subgap DOS andto accurately investigate the variations in device performance.The other two devices showed good device operation androbust bias stabilities, however, it was noteworthy that thedevice employing the 1 : 1 : 1 composition exhibited a lowerVTH, a higher msat, and larger degrees of PBTS/NBIS instabilitiesdue to a higher subgap DOS near the Ec and Ev than thedevice employing the 1 : 1 : 3 composition. These differenceswere suggested to originate from the different In/Zn ratios inthe ALD-IGZO channels. Furthermore, the negative DVTH underthe PB(T)S condition for the device employing the 1 : 1 : 1composition was also explained by the effects of the ALD cyclevariation. Consequently, the device employing the 1 : 1 : 3 com-position exhibited the most desirable device characteristicsfrom the viewpoint of excellent bias stability, and they werefound to be superior to those obtained from the conventionalsputter-deposited IGZO TFTs. Therefore, the compositioncontrol of the IGZO channel enabled by the ALD process is akey parameter to guarantee high performance and robuststability for future various practical applications using ALD-IGZOthin films.

Conflicts of interest

There are no conflicts to declare.

Acknowledgements

This work was partly supported by the Korea Evaluation Insti-tute of Industrial Technology through the Korean Government(10079974, Development of core technologies on materials,devices, and processes for TFT backplane and light emittingfrontplane with enhanced stretchability above 20%, with appli-cation to stretchable display) and by the Kyung Hee University-Samsung Electronics Research and Development Programentitled Flexible Flash Memory Device Technologies for Next-GenConsumer Electronics.

References

1 K. Nomura, H. Ohta, A. Takagi, T. Kamiya, M. Hirano andH. Hosono, Nature, 2004, 432, 488–492.

2 E. Fortunato, P. Barquinha and R. Martins, Adv. Mater.,2012, 24, 2945.

3 W. B. Yoo, C. K. Ha, J. W. Kwon, J. W. Jeon, H. J. Lee,S. Y. Kwon, W. S. Park, B. C. Kim, W. S. Shin and J. W. Kim,SID Symp. Dig. Tech. Papers, 2018, 49, 714–716.

4 A. J. Kronemeijer, I. Katsouras, P. Poodt, H. Akkerman,A. V. Breemen and G. Gelinck, Proc. 25th Int. WorkshopActive-Matrix Flatpanel Displays and Devices (AM-FPD),2018, 5-1.

5 K. Kaneko, N. Inoue, S. Saito, N. Furutake and Y. Hayashi,VLSI Symp. Dig. Tech. Papers, 2011, 120–121.

6 T. H. Yang, T. Y. Chen, N. T. Wu, Y. T. Chen and J. J. Huang,IEEE Trans. Electron Devices, 2017, 64, 1294–1299.

7 S. Knobelspies, B. Bierer, A. Daus, A. Takabayashi,G. A. Salvatore, G. Cantarella, A. O. Perez, J. Wollenstein,S. Palzer and G. Troster, Sensors, 2018, 18, 358.

8 Y. Fujimoto, M. Uenuma, Y. Ishikawa and Y. Uraoka, AIPAdv., 2015, 5, 097209.

9 B. Cui, L. Zeng, D. Keane, M. J. Bedzyk, D. B. Buchholz,R. P. H. Chang, X. Yu, J. Smith, T. J. Marks, Y. Xia, A. F.Facchetti, J. E. Medvedeva and M. Grayson, J. Phys. Chem. C,2016, 120, 7467–7475.

10 H. Chen, Y. Cao, J. Zhang and C. Zhou, Nat. Commun., 2014,5, 4097.

11 E. S. Hwang, J. S. Kim, S. M. Jeon, S. J. Lee, Y. J. Jang,D. Y. Cho and C. S. Hwang, Nanotechnology, 2018,29, 155203.

12 H. Yabuta, M. Sano, K. Abe, T. Aiba, T. Den, H. Kumomi,K. Nomura, T. Kamiya and H. Hosono, Appl. Phys. Lett.,2006, 89, 112123.

13 M. Leskela and M. Ritala, Angew. Chem., Int. Ed., 2003, 42,5548–5554.

14 V. Miikkulainen, M. Leskela, M. Ritala and R. L. Puurunen,J. Appl. Phys., 2013, 113, 021301.

15 A. Illiberi, B. Cobb, A. Sharma, T. Grehl, H. Brongersma,F. Roozeboom, G. Gelinck and P. Poodt, ACS Appl. Mater.Interfaces, 2015, 7, 3671–3675.

16 S. M. Yoon, N. J. Seong, K. J. Choi, G. H. Seo and W. C. Shin,ACS Appl. Mater. Interfaces, 2017, 9, 22676–22684.

17 S. J. Yoon, N. J. Seong, K. J. Choi, W. C. Shin and S. M. Yoon,RSC Adv., 2018, 8, 25014–25020.

18 T. Kamiya, K. Nomura and H. Hosono, J. Disp. Technol.,2009, 5, 468–483.

19 W. Korner, D. F. Urban and C. Elsasser, J. Appl. Phys., 2013,114, 163704.

20 W. H. Han and K. J. Kang, Phys. Rev. Appl., 2016, 6, 044011.21 T. Iwasaki, N. Itagaki, T. Den, H. Kumomi, K. Nomura,

T. Kamiya and H. Hosono, Appl. Phys. Lett., 2007, 90, 242114.22 P. Barquinha, L. Pereira, G. Gonçalves, R. Martins and

E. Fortunato, J. Electrochem. Soc., 2009, 156, H161–H168.23 H. S. Jeon, S. W. Na, M. R. Moon, D. G. Jung, H. S. Kim and

H. J. Lee, J. Electrochem. Soc., 2011, 158, H949–H954.24 M. H. Cho, H. J. Seol, H. C. Yang, P. S. Yun, J. U. Bae,

K. S. Park and J. K. Jeong, IEEE Electron Device Lett., 2018,39, 688.

25 M. H. Cho, M. J. Kim, H. J. Seul, P. S. Yun, J. U. Bae,K. S. Park and J. K. Jeong, J. Inf. Disp., 2018, DOI: 10.1080/15980316.2018.1540365.



26 H. K. Noh, K. J. Chang, B. K. Ryu and W. J. Lee, Phys. Rev. B:Condens. Matter Mater. Phys., 2011, 84, 115205.

27 T. Oh, Trans. Electr. Electron. Mater., 2017, 18, 21–24.28 T. Kamiya, K. Nomura and H. Hosono, Sci. Technol. Adv.

Mater., 2010, 11, 044305.29 H. J. Choi, J. M. Lee, H. Y. Bae, S. J. Choi, D. H. Kim and

D. M. Kim, IEEE Trans. Electron Devices, 2015, 62, 2689–2694.30 C. Chen, K. Abe, H. Kumomi and J. Kanicki, IEEE Trans.

Electron Devices, 2009, 56, 1177–1183.31 M. M. Billah, D. H. Chowdhury, M. Mativenga, J. G. Um,

R. K. Mruthyunjaya, G. N. Heiler, T. J. Tredwell and J. Jang,IEEE Electron Device Lett., 2016, 35, 735.

32 W. Meyer and H. Neldel, Z. Tech. Phys., 1937, 18, 588.33 R. Metselaar and G. Oversluizen, J. Solid State Chem., 1984,

55, 320–326.34 H. R. Im, H. S. Song, J. W. Jeong, Y. W. Song and Y. T. Hong,

Jpn. J. Appl. Phys., 2015, 54, 03CB03.

35 S. H. Jin, T. W. Kim, Y. G. Seol, M. Mativenga and J. Jang,IEEE Electron Device Lett., 2014, 35, 560.

36 Y. H. Chang, M. J. Yu, R. P. Lin, C. P. Hsu and T. H. Hou,Appl. Phys. Lett., 2016, 108, 033502.

37 J. Y. Huh, J. H. Jeon, H. H. Choe, K. W. Lee, J. H. Seo,M. K. Ryu, S. H. K. Park, C. S. Hwang and W. S. Cheong, ThinSolid Films, 2011, 519, 6868–6871.

38 S. J. Choi, J. T. Jang, H. R. Kang, J. H. Baeck, J. U.Bae, K. S. Park, S. Y. Yoon, I. B. Kang, D. M. Kim,S. J. Choi, Y. S. Kim and S. Oh, IEEE Electron Device Lett.,2017, 38, 580.

39 H. C. Oh, S. M. Yoon, M. K. Ryu, C. S. Hwang, S. H. Yang andS. H. K. Park, Appl. Phys. Lett., 2010, 97, 183502.

40 W. S. Cheong, J. H. Park and J. K. Shin, ETRI J., 2012,34, 966.

41 K. A. Kim, M. J. Park, W. H. Lee and S. M. Yoon, J. Appl.Phys., 2015, 118, 234504.


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