144
AN ABSTRACT OF THE THESIS OF
Jaana Saranya Rajachidambaram for the degree of Master of Science in Chemical
Engineering presented on December 21, 2011
Title: Evaluation of Amorphous Oxide Semiconductors for Thin Film Transistors (TFTs)
and Resistive Random Access Memory (RRAM) Applications
Abstract approved:
Gregory S. Herman
ABSTRACT
Thin-film transistors (TFTs) are primarily used as a switching element in liquid crystal
displays. Currently, amorphous silicon is the dominant TFT technology for displays, but
higher performance TFTs will become necessary to enable ultra-definition resolution
high-frequency large-area displays. Amorphous zinc tin oxide (ZTO) TFTs were
fabricated by RF magnetron sputter deposition. In this study, the effect of both deposition
and post annealing conditions have been evaluated in regards to film structure,
composition, surface contamination, and device performance. Both the variation of
oxygen partial pressure during deposition and the temperature of the post-deposition
annealing were found to have a significant impact on TFT properties. X-ray diffraction
data indicated that the ZTO films remain amorphous even after annealing to 600° C.
Rutherford backscattering spectrometry indicated that the Zn:Sn ratio of the films was
~1.7:1 which is slightly tin rich compared to the sputter target composition. X-ray
photoelectron spectroscopy data indicated that the films had significant surface
contamination and that the Zn:Sn ratios changed depending on sample annealing
conditions. Electrical characterization of ZTO films using TFT test structures indicated
that mobilities as high as 17 cm2 V-1 s-1 could be obtained for depletion mode devices. It
was determined that the electrical properties of ZTO films can be precisely controlled by
varying the deposition conditions and annealing temperature. It was found that the ZTO
electrical properties could be controlled where insulating, semiconducting and conducting
films could be prepared. This precise control of electrical properties allowed us to
incorporate sputter deposited ZTO films into resistive random access memory (RRAM)
devices. RRAM are two terminal nonvolatile data memory devices that are very
promising for the replacement of silicon-based Flash. These devices exhibited resistive
switching between high-resistance states to low-resistance states and low-resistance states
to high-resistance states depending on polarity of applied voltages and current
compliance settings. The device switching was fundamentally related to the defect states
and material properties of metal and insulator layers, and their interfaces in the metal-
insulator-metal (MIM) structure.
©Copyright by Jaana Saranya Rajachidambaram
December 21, 2011
All Rights Reserved
Evaluation of Amorphous Oxide Semiconductors for Thin Film Transistors (TFTs) and
Resistive Random Access Memory (RRAM) Applications
by
Jaana Saranya Rajachidambaram
A THESIS
submitted to
Oregon State University
in partial fulfillment of
the requirements for the
degree of
Master of Science
Presented December 21, 2011
Commencement June 2012
Master of Science thesis of Jaana Saranya Rajachidambaram presented on December 21,2011
APPROVED :
Major Professor, representing Chemical Engineering
Head of School of Chemical, Biological & Environmental Engineering
Dean of the Graduate School
I understand that my thesis will become a part of the permanent collection of Oregon State University libraries. My signature below authorizes release of my thesis to any reader upon request
Jaana Saranya Rajachidambaram, Author
ACKNOWLEDGEMENTS
First and foremost, I would like to thank my mom, Chithra and dad,
Rajachidambaram for their sincere prayers, enormous concern and support throughout
my graduate life. I would also like to thank my sweet sister Priya and my brother-in-law,
Gopi who encouraged me towards pursuing MS degree in the US and took care of my
initial living expenses. I should thank their new born baby Akanshaa who kept me
cheerful and enthusiastic during the month of my MS defense.
I would like to specifically thank my advisor, Dr. Gregory S. Herman for his
guidance, patience, encouragement and funding to carry out my research and I am happy
to have been one of his first two students in the research group. Though I had a lot of
responsibilities, I enjoyed learning new things while being in the research group.
I would like to express my heartfelt thanks to scientists from Pacific Northwest
National Laboratory (PNNL), Dr. Thevuthasan for offering me an Alternate Sponsored
Fellowship to perform a portion of this research work and Dr. Nachimuthu, Dr.Varga, Dr.
Shutthanandan and S.Sanghavi for their help with data acquisition and analysis. I would
also like to thank Dr. S.L. Golledge from CAMCOR for his guidance with acquiring
SIMS data.
I would like to recognize all my research group members including Brendan
Flynn and Richard Oleksak who took time to review my thesis documents and Eric
Hostetler, Udit Suri, Lane Porth, Charith Abeywarna, who have helped me in
accomplishing research tasks.
Special thanks to Chris Tasker for his help with using research tools in the OSU
cleanroom. A warm thanks to Santosh Murali and Dr.Conley for their assistance with the
memristor project. I would like to thank Dr.Chang and Dr.Han for their participation and
inputs in the weekly memristor meeting.
I am very grateful to Dr.Williamson and Janet Mosley for the scholarship offer
and their guidance.
The substrates for TFTs and RRAMs were provided by Hewlett-Packard (HP) and Sharp
Laboratories of America (SLA).
The research was performed using facilities at the Microproducts Breakthrough
Institute at Oregon State University and at the Environmental Molecular Sciences
Laboratory (EMSL), a national scientific user facility at Pacific Northwest National
Laboratory (PNNL) sponsored by the U.S Department of Energy. JSR thanks PNNL for
providing an Alternate Sponsored Fellowship during a portion of these studies.
The project was funded by the Oregon Nanoscience and Microtechnologies
Institute (ONAMI) and the Office of Naval Research (ONR) under contract number
200CAR262.
TABLE OF CONTENTS
Page
CHAPTER 1 - INTRODUCTION TO OXIDE MATERIALS AND DEVICES…. 1
1.1 Introduction on Thin-Film Transistors………………………………… 1
1.2 TFT device structure and Operation…………………………………… 2
1.3 Transparent Amorphous Oxide Semiconductors………………………. 8
1.3.1 Indium Gallium Zinc Oxide (IGZO)………………………… 10
1.3.2 Zinc Tin Oxide (ZTO)……………………………………….. 12
1.3.3 Zinc Indium Oxide (ZIO)……………………………………. 14
1.3.4 Indium Gallium Oxide (IGO)………………………………... 15
1.4 Introduction to Resistive Random Access Memory…………………….17
1.5 RRAM device structure and operation………………………………… 18
1.6 Oxide RRAM Overview……………………………………………….. 29
1.7 References ………………………………………………………………32
CHAPTER 2 - EXPERIMENTAL TECHNIQUES……………………………….. 39
2.1 Thin Film Deposition Techniques………………………………………39
2.1.1 Radio-Frequency Magnetron Sputtering……………………...39
2.1.2 Thermal Evaporator…………………………………………...42
2.2 Lithography…………………………………………………………….. 43
TABLE OF CONTENTS (Continued)
Page
2.2.1 Shadow Mask Lithography……………………………………43
2.2.2 Photolithography………………………………………………44
2.3 Post-deposition annealing……………………………………………….46
2.4 Thin film characterization techniques………………………………….. 47
2.4.1Spectroscopic Ellipsometry……………………………………47
2.4.2 Scanning Electron Microscopy………………………………. 51
2.4.3 X-ray Diffraction…………………………………………….. 52
2.4.4 X-ray Photoelectron Spectroscopy……………………………53
2.4.5 Rutherford Backscattering Spectrometry……………………..56
2.5 Electrical characterization using semiconductor parameter analyzer….. 59
2.6 References………………………………………………………………60
CHAPTER 3 – INVESTIGATION OF AMORPHOUS ZINC TIN OXIDE FILMS FOR THIN FILM TRANSISTOR APPLICATIONS……………………. 62
3.1 Introduction…………………………………………………………….. 62
3.2 Experimental details…………………………………………………….64
3.3 Results and discussion…………………………………………………. 67
3.4 Conclusion………………………………………………………………82
TABLE OF CONTENTS (Continued)
Page
3.5 References ………………………………………………………………83
CHAPTER 4 – BIPOLAR RESISTIVE SWITCHING IN SPUTTER DEPOSITED AMORPHOUS ZINC TIN OXIDE DEVICES…………………………………… 86
4.1 Introduction……………………………………………………………. 86
4.2 Experiments…………………………………………………………… 88
4.3 Results and discussion………………………………………………… 89
4.4 Conclusion………………………………………………………………102
4.5 References ………………………………………………………………103
CHAPTER 5 - CONCLUSIONS AND RECOMMENDATIONS FOR FUTURE WORK……………………………………………………………………………... 107
5.1 Conclusions……………………………………………………………. 107
5.2 Recommendations for future work…………………………………….. 108
Bibliography……………………………………………………………………….. 109
Appendix: Sol-gel chemistry for zinc tin oxide (ZTO)……………………………. 124
LIST OF FIGURES
Figure Page
Figure 1.1: Four basic TFT layouts: (A) staggered bottom gate, (B) staggered top gate, (C) coplanar bottom gate, (D) coplanar top gate………………………………3
Figure 1.2: Energy band diagram for n-type accumulation mode TFTs.(A) at zero bias (B) negative voltage is applied to the gate (C) positive voltage is applied to the gate………………………………………………………………………………….4
Figure 1.3: Plot of log(ID) and log(IG) versus VGS (V) to illustrate the on-off ratio and turn-on voltage…………………………………………………………………6
Figure 1.4: Plot of VGS (V) versus µ (cm2/V·s)…………………………………….7
Figure 1.5: Atomic orbital overlap (shaded) (A) Crystalline covalent semiconductors, (B) Amorphous covalent semiconductors, (C) Crystalline oxide semiconductors and (D) Amorphous oxide semiconductors……………………… 9
Figure 1.6: A portion of the periodic table for selecting AOS materials excluding expensive (shaded white) and toxic (shaded red) elements. Gallium and indium are also expensive……………………………………………………………………...10
Figure 1.7: Side view and top view for RRAM device cross bar structure………..19
Figure 1.8: AUTOCAD photomask design for 10000 um2 cross bar structure……20
Figure 1.9: Bipolar resistive switching in RRAM…………………………………21
Figure 1.10: Unipolar resistive switching in RRAM………………………………22
Figure 1.11: Gentle forming process – soft breakdown of oxide material…………23
Figure 1.12: Forming, SET and RESET operation in RRAM leading to conductive filament formation and rupture…………………………………………………...... 26
LIST OF FIGURES(Continued)
Figure Page
Figure 1.13: Resistive switching due to redox reactions at the metal electrode/oxide Interface…………………………………………………………… 28
Figure 2.1: Schematic of sputter deposition system……………………………….. 41
Figure 2.2: Schematic of thermal evaporator system. ………………………………43
Figure 2.3: Schematic of the dark field and bright field photomask sets used in this study for fabrication of RRAM cross bar structures : (A) Bottom electrode mask, (B) : Pad mask and (C) Top electrode mask………………………………… 46
Figure 2.4: Ellipsometer for measurement of thickness and optical constants…….. 47
Figure 2.5: Two kinds of MSE minima called local minima and global minima...... 50
Figure 2.6: SEM image of platinum bottom electrode of ZTO RRAM cross bar structure…………………………………………………………………………… 52
Figure 2.7: Principle of X-ray diffraction………………………………………….. 52
Figure 2.8: Principle of XPS (left) and ejection of a photoelectron from inner shell of an atom (right)…………………………………………………………………... 54
Figure 2.9: XPS spectra for Zn 2p3/2 with binding energy and photoelectron intensity as X and Yaxis…………………………………………………………… 55
Figure 2.10: Principle of RBS……………………………………………………… 56 Figure 2.11: RBS spectra of ZTO deposited on SiO2/Si substrates with channel number and backscattering yield as X and Y axis…………………………………. 58
Figure 3.1: XRD spectra for ZTO thin films as-deposited and annealed to 600° C and 650° C…………………………………………………………………… . .. 68
Figure 3.2: SEM image of as-deposited ZTO film taken at 65,000x magnification. 69
LIST OF FIGURES (Continued)
Figure Page
Figure 3.3: XPS of Zn 2p3/2 from ZTO films deposited with 1 SCCM of oxygen and annealed to 300 °C and 600 °C………………………………………………………. 70
Figure 3.4: XPS of Sn 3d5/2 from ZTO films deposited with 1 SCCM of oxygen and annealed to 300 °C and 600 °C………………………………………………………. 71
Figure 3.5: XPS of O 1s from ZTO films deposited with 1 SCCM of oxygen and annealed to 300 °C and 600 °C………………………………………………………. 72
Figure 3.6: XPS of C 1s from ZTO films as-deposited with 1 SCCM of oxygen and annealed to 300 °C and 600 °C……………………………………………………… 73
Figure 3.7: XPS sputter depth profile of ZTO film deposited with 1 SCCM of oxygen……………………………………………………………………………….. 75
Figure 3.8: RBS spectra from an as-deposited ZTO film with 0 SCCM of oxygen…. 76
Figure 3.9: TFT transfer characteristics with as-deposited ZTO channel layer with 0 SCCM, 1 SCCM, 2 SCCM of oxygen………………………………………………. 78
Figure 3.10: TFT transfer characteristics with ZTO channel layer deposited with 0 SCCM, 1 SCCM, 2 SCCM of oxygen and annealed to 300 °C……………………… 78
Figure 3.11: TFT transfer characteristics with ZTO channel layer deposited with 0 SCCM, 1 SCCM, 2 SCCM of oxygen and annealed to 600 °C……………………… 79
Figure 3.12: Change in average channel mobility as a function of oxygen flow rate and post-annealing conditions……………………………………………………….. 81
Figure 4.1: RBS data from sputter-deposited ZTO films. The inset shows the XRD data from the ZTO films after annealing to 600 °C…………………………………. 90
Figure 4.2: ToF-SIMS depth profile on Al/ZTO/Pt structure………………………... 92
Figure 4.3: Bipolar resistive switching in ZTO RRAM crossbar device; inset shows an image of device structure………………………………………………... 94
LIST OF FIGURES(Continued)
Figure Page
Figure 4.4: Bipolar resistive switching in ZTO RRAM crossbar device; inset shows a schematic of the device stack and dependence of RLRS and RHRS on cell area………………………………………………………………………................... 95
Figure 4.5: Ohmic conduction at low voltage regions of LRS and HRS (V < V
RESET) and the insets show plots for Schottky emission and Poole – Frenkel emission for the high voltage regions (V > VRESET) of HRS……………………... 97
Figure 4.6: Statistical distribution of SET and RESET voltages for 30 cycles where µ denotes mean and σ denotes standard deviation………………………... 99
Figure 4.7: RLRS and RHRS over 30 resistive switching cycles………………………. 100
Figure 4.8: Retention characteristics of ZTO memory device taken for a read voltage of 0.1 V…………………………………………………………………..….. 100
LIST OF TABLES
Table page
Table 3.1: XPS table of binding energies for Zn 2p3/2, Sn 3d5/2, O 1s and C 1s for ZTO films with different deposition conditions and annealing temperature……… … 74
Table 3.2: XPS table of atomic ratios for as-deposited ZTO thin films……………... 75
Table 3.3: Fitting parameters for RBS spectra obtained using simulation……………77
Table 4.1: Experimental values for optical dielectric constant (εr) for ZTO calculated using slopes (S1 and S2) of Schottky and Poole-Frenkel conduction mechanism at the high voltage regions (V > VRESET) of HRS. The optical dielectric constant εr,o for ZTO was determined to be ~ 4 from spectroscopic ellipsometry….. 97
1
CHAPTER 1 - INTRODUCTION TO OXIDE MATERIALS AND DEVICES
This chapter provides an overview of amorphous oxide semiconductor materials and
device structures that have been reported in the literature. Information on the principle
concepts of thin film transistors (TFT) and resistive random access memory (RRAM) will
also be covered.
1.1 Introduction on Thin-Film Transistors
The first metal-oxide-semiconductor field-effect transistor (MOSFET) was
fabricated in 1960 on a silicon substrate using SiO2 as the gate dielectric. Several of these
MOSFETs were combined to fabricate integrated circuits by Texas Instruments and
Fairchild Semiconductor [1, 2]. In 1963, the invention of complementary MOS (CMOS)
was a major breakthrough where both n-channel and p-channel MOSFET were integrated
into circuits with the advantage that negligible standby power dissipation for CMOS as
opposed to PMOS or NMOS. Thin film transistors (TFT) are a special kind of field effect
transistor where the gate insulator, channel layer and the metal contacts are deposited as
thin films. TFTs have been studied since their invention in 1930 [3-6] and used for a
broad range of applications, but with liquid crystal displays being the dominant
application and amorphous-silicon (a-Si) being the dominant TFT technology.
A recent advance for TFTs was realized in 1996 when oxide semiconductors were
first used as the channel material for TFTs. These materials initially were focused on tin
oxide (SnO2) as the semiconductor where antimony (Sb) was used as a dopant [7].
However, in 2003 polycrystalline ZnO was used as the channel material [8] and these
2
studies indicated that ZnO TFTs had significantly better performance than a-Si based
TFTs [8]. Considerable advancements have been made for ZnO based TFTs including
enhanced electrical properties with low processing temperatures [9]. More recently a
general class of materials commonly termed multi-component oxide based
semiconductors has been studied which includes, indium gallium zinc oxide (IGZO) [10],
zinc tin oxide (ZTO) [11], indium zinc oxide (IZO) [12], and indium gallium oxide (IGO)
[13]. It has been demonstrated that these amorphous materials have relatively high
electron mobilities despite being amorphous [14]. For the research in this thesis the focus
has been on ZTO, which is an amorphous oxide semiconductor.
1.2 TFT Device structure and Operation
Figure 1.1 shows the schematic for four basic TFT layouts: (A) staggered
bottom-gate, (B) co-planar bottom-gate, (C) staggered top-gate, and (D) co-planar top-
gate [15]. The staggered TFT device structures have the source/drain contacts and gate
contacts placed on opposite sides of the channel-insulator interface and in co-planar TFT
device structures the source/drain contacts and gate contacts are placed on the same side
of the channel-insulator interface. TFTs are also classified as top and bottom gate based
on the placement of the gate electrode in the device stack. The top gate TFTs have the
gate electrode placed on the top of the channel layer and bottom gate TFTs have the gate
electrode beneath the channel layer. There are several advantages and disadvantages
associated with each structure. In this study, the staggered bottom gate structure is
employed.
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5
where Cox is the capacitance of the gate oxide, W and L are the width and length of the
channel, VON is the voltage at which electrons start to flow through the channel and
current conduction begins, µ is the channel mobility, VGS and VDS are the gate-to-source
and drain-to-source voltages, respectively[16].
When a large drain voltage is applied, electrons are depleted in the channel layer
near the drain electrode and ID becomes independent of drain voltage.
= 12μ [( − ) ], = ( − ) A derivation of the ideal square law model is available elsewhere [16]. All the
devices were tested using a probe station with an Agilent 4155 semiconductor parameter
analyzer. The I-V curves were obtained using the double sweep mode with gate voltage
ranging from -20 to 20 V with drain voltage fixed at 1V.
Turn-On Voltage
The turn-on voltage (Von) is the gate voltage at which the drain current increases
due to current conduction in the channel. The plot of log (ID) versus VGS shown in Figure
1.3 is used to determine the value for Von where an arrow indicates the voltage at which
the drain current starts to increase.
6
Figure 1.3: Plot of log (ID) and log (IG) versus VGS (V) to illustrate the on-off ratio and turn-on voltage.
Drain current ON to OFF Ratio
The drain current on-to-off ratio is defined to be the ratio of drain to source
current from the “on” state to the “off” state obtained from the plot of log (ID) versus VGS
shown in Figure 1.3. Typically, on-off ratios are considered acceptable when their values
are greater than 106 [17]. On-off ratios are useful in accessing how well a device will
work as a switch.
Mobility
Mobility is a measure of the transport of carriers in the channel. A higher mobility
contributes to a faster switching time, which is indicative of the time the device takes to
switch between the on and off states. Two types of mobility calculated in this study that
-12-11-10-9-8-7-6-5-4-3-2
-12-11-10
-9-8-7-6-5-4-3-2
-30 -20 -10 0 10 20 30
log
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log
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))
VGS (V)
Von
ON-OFF ratio
7
depends on VGS are incremental mobility (µinc), which is the increase in mobility due to
carriers added to the channel as the gate voltage differentially increases in magnitude, and
average mobility (µavg), which is the average mobility of all the carriers present in the
channel [8]. A plot of VGS (V) versus µ (cm2/V·s) is shown in Figure 1.4.
Figure 1.4: Plot of VGS (V) versus µ (cm2/V·s).
The incremental mobility is calculated by
μ ( ) = ( ) where G’
CH is the differential channel conductance as a function of gate voltage which
can be calculated from channel conductance data and is equal to ∆GCH / ∆ VGS , W/L are
0
5
10
15
20
25
30
-30 -20 -10 0 10 20 30
µ (c
m2 /
V. s)
VGS (V)
µincµavg
8
the width to length ratio of the channel layer between the source and drain, and Cins is the
capacitance of the insulator.
The average channel mobility is calculated by
μ = ( ) [ − ] where GCH is the channel conductance as a function of gate voltage VGS, and VON is the
voltage at which the device switches on.
1.3 Transparent Amorphous Oxide Semiconductors
In 1996 amorphous oxide semiconductor (AOS) materials were explored for their
application as transparent conducting oxides (TCO) [18]. More recently these strategies
have been applied to TFT applications as well [19]. For these applications it was found
that AOS materials were promising due to their low processing temperature, lack of grain
boundaries and smoothness of the surface. These metal oxides are composed of spatially
spread s-orbitals of heavy metal cations (HMCs) with an electronic configuration of (n-
1)d10ns0 [18,20]. The magnitude of atomic orbital overlap is proportional to electron
mobility in the material. When compared to crystalline Si shown in Figure 1.5A, a-Si is
greatly affected by structural randomness as their carrier transport paths are composed of
strongly directional sp3 orbitals as shown in Figure 1.5B. In case of crystalline oxide
semiconductors as shown in Figure 1.5C, the overlap of spherical s metal orbitals are
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11
TFTs had good performance with an on-off ratio of 106 and a field effect mobility of 80
cm2/V·s. In 2004, Nomura et al. demonstrated the room temperature fabrication of
amorphous IGZO TFTs on flexible polymer substrates. The IGZO was deposited using
pulsed laser deposition (PLD). The TFTs displayed good performance with saturation
mobilities of ~ 6 to 9 cm2/V·s, an on-to-off ratio of 103 before and after bending [23].
Following this initial work, several research groups focused on improving the IGZO TFT
performance. For example, in 2011 Liu et al. demonstrated that doping nitrogen into the
a-IGZO film during deposition improved performance of the IGZO TFT [25]. These
TFTs were fabricated with an inverted staggered bottom gate structure where the active
channel layer was formed using DC reactive sputtering with a power of 100 W and
deposition was performed with Ar and N2 gas flow in the deposition chamber. It was
proposed that the nitrogen was incorporated into the a-IGZO film, and substituted
nitrogen atoms for inactive oxygen atoms. The inactive oxygen atoms can react with
ambient air and form oxygen vacancies, which significantly affect the electrical and
ambient stability of IGZO TFTs. Experimental results show that TFT device parameters
including threshold voltage, sub-threshold swing, and carrier mobility were enhanced as a
result of nitrogen doping. The devices have also proven to be electrically stable under
applied gate bias stress and were stable in ambient conditions. This was attributed to
reducing the oxygen desorption effect caused by interaction of inactive oxygen atoms of
IGZO with the atmosphere through the incorporation of nitrogen into film during
deposition [25].
12
1.3.2 Zinc Tin Oxide (ZTO)
Zinc tin oxide, also called zinc stannate, is a wide band gap (3.3-3.9 eV) n-type
semiconductor material, which is also commonly investigated for TFT applications. The
ZTO stoichiometry can be described generally as (ZnO)x(SnO2)1-x with 0 < X < 1. The
two crystalline forms of ZTO that have been studied in the literature are ilmenite ZnSnO3
(x=0.5) and cubic spinel Zn2SnO4 (x = 0.66), with the ilmenite structure exhibiting more
thermal stability [11, 26]. ZTO has many attractive attributes, such as its resistance to
chemical etching [27], and physical robustness to scratching. Moreover, Zn and Sn are
relatively low cost and abundantly available. Besides TFT applications, ZTO has also
been used for gas sensor [28] and solar cell applications [29].
It has been demonstrated that ZTO films can be formed by many deposition methods
including vacuum based techniques like RF magnetron sputter deposition [11,21],
chemical vapor deposition [30] and solution based techniques including spin coating [31-
33], ink jet printing [34,35] and dip coating [36].
In 2005, Chiang et al. demonstrated n-type bottom gate TFTs with ZTO as a channel
layer [11]. These TFTs were fabricated using glass substrates coated with a 200 nm thick
gate electrode composed of sputter deposited ITO and a 220 nm thick gate insulator
composed of alternating AlOx and TiOx (ATO) deposited by atomic layer deposition. The
sputter deposition of the ZTO channel layer and ITO source/drain was done with a
90%/10% Ar/O2 mix and 100% Ar at a substrate temperature of ~175 °C. The ZTO films
were deposited from a target containing 1:1 and 2:1 molar ratios of ZnO:SnO2 . The
13
devices were furnace annealed for 1 hour in air at 300 °C and 600 °C following the ZTO
deposition. The XRD results indicated the ZTO films were amorphous up to 600 °C and
polycrystalline above 650 °C and the entire device structure with ZTO as channel layer
had high transparency with ~ 84 % optical transmittance. The TFTs annealed at 300 °C
had field effect mobilities of 5 to15 cm2/V·s and a turn-on voltage of 0 to15 V whereas
the TFTs annealed at 600 °C had field effect mobilities of 20 to 50 cm2/V·s and a turn-on
voltage of -5 to 5 V [11].
Following the successful fabrication of TFTs using ZTO as a channel layer, research has
focused on improving the electrical stability of ZTO based TFTs. For example, in 2011
Avis et al. investigated the stability of unpassivated ZTO based TFTs to assess the impact
of air exposure and bias stress [37]. The ZTO thin film was formed with bottom gate,
bottom contact structure on glass substrate with 200 nm of SiO2 buffer layer deposited on
top of it using plasma-enhanced chemical vapor deposition (PECVD) [37]. The
source/drain and gate electrodes were formed using sputter deposited IZO (40 nm). The
gate dielectric was a SiO2 layer (150 nm) and the ZTO was deposited on top of the TFT
backplane by spin coating from a solution containing ZnCl2 and SnCl2 dissolved in
acetonitrile and ethylene glycol with Zn : Sn ratio of 1:1, 2:1, and 1:2. Annealing was
performed at 500 °C for 1 hr in air and resulting thickness of ZTO was 20 nm. XRD
analysis showed that the films were amorphous at 500 °C. Electrical measurements
showed that the best values for linear mobility and on-off ratio were 6.77 cm2/V·s and
106 respectively, which were obtained for TFTs fabricated with ZTO channel layer
containing a Zn : Sn ratio of 1:1.The impact of air exposure on TFTs increased the off-
14
current by approximately six orders of magnitude. The impact of positive gate bias stress
revealed a large threshold voltage (Vth) shift. On the other hand, the impact of negative
gate bias stress revealed no marked Vth shifts. This study illustrated the necessity of
having a passivation layer above the ZTO channel layer for stable TFT performance [37].
1.3.3 Zinc Indium Oxide (ZIO)
Amorphous zinc indium oxide is best known for its excellent optical transparency,
high electrical conductivity, thermal and electrical stability and film smoothness for TCO
applications [38-40].
In 2005 Dehuff et al. reported the fabrication of n-type bottom gate TFTs with
ZIO as a channel layer [40]. These TFTs were fabricated using glass substrates coated
with a 200 nm thick gate electrode composed of sputter deposited ITO and a 220 nm
thick gate insulator composed of alternating layers of AlOx and TiOx (ATO) deposited by
atomic layer deposition. The sputter deposition of the channel layer ZIO was done at
room temperature using a target made from 2:1 molar ratio of ZnO:In2O3. Process
parameters were varied, including oxygen flow rate and target-to-substrate distance
during ZIO channel layer deposition. The devices were furnace annealed in air for 1 hour
at 300 °C and 600 °C following the ITO source/drain deposition. The XRD results
indicated the ZIO films to be amorphous up to 500 °C and polycrystalline at 600 °C and
the devices had high transparency with ~ 85 % optical transmission. The TFTs
performance varied significantly with annealing temperature, where TFTs annealed at
600 ° C operated in depletion mode with threshold voltages between -20 and -10 V,
15
incremental channel mobility of 45 - 55cm2/V.s, and drain current on-off ratios of 106. In
contrast the TFTs annealed at 300 ° C operated in enhancement mode with threshold
voltages between 0 and 10 V, incremental channel mobility of 10 – 30 cm2/V.s, and drain
current on-off ratios of 106 [40].
In 2011 Han et al., evaluated the effect of ozone annealing in comparison to air
annealing on IZO TFT performance for solution processed IZO [41]. Bottom gate TFTs
were fabricated on boron doped silicon substrates with a 100 nm silicon oxide gate
insulator, and a 500 nm gold layer deposited on the back side of the substrate for gate
contact. The IZO thin film was spin coated at 300 rpm for 30 seconds using a precursor
prepared by dissolving InCl3 and ZnCl2 precursors in acetonitrile solvent. The films were
annealed for 2 hours in air and O2/O3 at temperatures between 280 °C and 500 °C. X-ray
photoelectron spectroscopy and atomic force microscopy studies were performed to study
the IZO film stoichiometry and morphology. TFT electrical device characterization
indicated that the TFTs prepared by annealing in ozone at 300 °C had a field effect
mobility of 0.94 cm2/V·s which was 30 times higher than the value of field effect
mobility by annealing in air at 300 °C. In this study it was shown that high mobility TFTs
could be obtained at low processing temperature using an ozone anneal [41].
1.3.4 Indium Gallium Oxide (IGO)
IGO has mainly been studied for application as a TCO [42, 43]. IGO is an
alternative amorphous oxide n-type semiconductor material with a wide band gap of 3.3-
3.4 eV [44]. In 2006, Chiang et al. fabricated IGO TFTs with staggered bottom gate
16
structure employing p-doped Si substrate with 100 nm SiO2 gate dielectric. The IGO
channel material and ITO source/drain material were RF magnetron sputter deposited and
patterned using shadow masks [44]. The sputter deposition was done using two ceramic
targets with different In/Ga ratios (InGaO3 and In1.33Ga0.67O3). Following the channel
layer deposition, the devices were furnace annealed in air at temperatures between 200 °C
and 800 °C. The effect of channel layer stoichiometry, post-deposition annealing, and
oxygen partial pressure on TFT electrical performance was studied and it was found that
oxygen partial pressure variation had a significant impact on the TFT performance. The
µinc of the devices increased with decreasing oxygen partial pressure whereas the VON
decreased (negative value) with decrease in oxygen partial pressure. The highest value
for the µinc was found to be 27 cm2/ V·s, with Von and on-off ratio measured as -14 V and
106 respectively for TFTs annealed at 600 °C. Additionally the µinc was found to be 19
cm2/ V·s with VON of 2 V for TFTs annealed at 200 °C [44].
In 2010, further research was done by Goncalves et al. to understand the
structural, morphological, optical and electrical properties of IGO thin films as a function
of oxygen partial pressure [45]. IGO films were deposited onto soda lime glass using RF
magnetron sputter deposition at room temperature using a 3 in. ceramic target made with
2:1 ratio of In2O3:Ga2O3. IGO films at room temperature and those annealed to 150 °C
were employed to fabricate staggered bottom gate TFTs .The XRD results indicated that
the films tend to be crystalline when deposited in the absence of oxygen and were
amorphous when deposited in the presence of oxygen. Film morphology studies using
AFM indicated that amorphous films were smoother (RMS = 1.2 nm) when compared to
17
the polycrystalline films (RMS = 11.5 nm) and the films deposited in the presence of
oxygen had 80% transparency when compared to films deposited in absence of oxygen
had less than 10% transparency. High performance TFTs were produced at room
temperature with high saturation mobility of 43 cm2/ V·s and on-off ratio of 108 [45].
1.4 Introduction to Resistive Random Access Memory
Resistive random access memory (RRAM) is a modern semiconductor non-
volatile memory which has attracted significant attention in the electronic industry due to
its simple structure, low power consumption, and fast operation speed [46]. When
compared to flash memory applications, RRAM requires a lower operating voltage
compared to several other competing technologies and is suitable for low power
applications. Since the early 1960s there was significant effort to study resistance
switching effects in thin films. In 1962, negative resistance was first observed in oxide
thin films including SiOx, Ta2O5, ZrO2, Al2O3 and TiO2 confirmed from current-voltage
characteristics [47]. A peak-to-valley ratio of 30:1 and a switching time of < 0.5 µs was
achieved but the mechanism behind negative resistance was unclear. Later in 1967 a
memory device was proposed with SiO thin films where change in resistance state was
reported to be due to electron penetration in electron-beam scanned areas, causing an
increase in the generated current within the device [48]. A review of switching observed
in various metal oxide films was put forward in 1970 which included discussion on
charge transport in insulators, observations and models of the forming process, and
differential negative resistance and its possible mechanisms [49].
18
A major advancement took place with the integration of RRAM into a 64 bit non-
volatile memory based on perovskite oxide, PCMO (Pr0.7Ca0.3MnO3), as a switching
element [50]. Beginning in 2004, a significant amount of research has focused on the
fabrication of RRAM based devices using binary transition metal oxides (TMO),
including NiO[51],TiO2[52], and CoO[53] as switching materials. These compounds
were found to have stable resistive switching where the best performance was found for
NiO films which had the lowest programming voltage and current. For the past seven
years, the research based on RRAM with oxides as the switching material has been
accelerated with non-volatile resistive switching observed in doped perovskite oxide
SrZrO3 [54], Al2O3[55], ZnO[56], ZrO2[57], HfOX[58], CuXO,[59] IGZO[60] and gallium
zinc oxide (GZO) [61].
To investigate the underlying mechanism of resistive switching a variety of
characterization techniques were used including conductive atomic force microscopy (C-
AFM), Auger electron spectroscopy (AES), and transmission electron microscopy
(TEM). In addition, analysis of I-V characteristics including change in resistance with
RRAM cell area have been performed.
1.5 RRAM device structure and operation
RRAM devices have a capacitor like structure in which the switching material is
sandwiched between two metal electrodes as shown in Figure 1.7. In 2008, HP
announced the development of a new circuit element termed the memristor [62]. The
memristor devices in this study have metal- insulator- metal (MIM) structures arranged in
a
ar
F
F
um
1
w
m
or
crossbar fas
rrays of elec
igure 1.7: Si
or our studie
m2, 2500 um
0000 um2 cr
while the top
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r negative vo
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ctrodes. This
ide view and
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m2 and 10000
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0 um2). A sc
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m electrode w
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m for the stru
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ed in four dif
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e 1.8. The bo
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round and th
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bar structure
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AD photoma
ottom electro
in this study
he top electro
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.
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ask design fo
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19
400
or
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igure 1.8: A
For th
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AUTOCAD p
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20
itially
e in
nce
age at
ve
tance
e
RS) of
21
the device and RON refers to low resistance state (LRS) of the device after the SET
process.
Figure 1.9: Bipolar resistive switching in RRAM.
There are two kinds of resistive switching, bipolar and unipolar. In bipolar
resistive switching the SET and RESET operation of the device depends on the polarity
of the applied voltage as described earlier whereas in unipolar resistive switching the
SET and RESET operation of the device depends on the amplitude of the applied voltage
[63]. Figures 1.9 and 1.10 depict the bipolar and unipolar behavior of the RRAM devices.
-3 -2 -1 0 1 2 3
Curr
ent (
I)
Voltage(V)
Bipolar resistive switching
CC
R ON
Pre - SET current
SET Voltage
Pre - RESET current
Post - RESET current
SET 10-2
10-3
10-4
10-5
10-6
10-1
100
R OFF
RESET Voltage
RESET
10-7
10-8
22
Figure 1.10: Unipolar resistive switching in RRAM.
Forming process
Forming refers to the soft breakdown of the oxide material (insulator) by applying
an external voltage on the top electrode, Al in this case. Figure 1.11 shows the effect of
slowly increasing the current for a device initially in the HRS. It has been proposed that
the forming process gives rise to the formation of a filamentary conductive path that
connects the top and bottom electrodes. This filamentary formation may possibly be due
to the application of bias voltage polarity on the top electrode. This could cause the
migration of oxygen ions (O2-) in one direction, or the migration of oxygen vacancies
-3.5 -3 -2.5 -2 -1.5 -1 -0.5 0
Curr
ent (
I)
Voltage (V)
Unipolar resistive switching
RESET SET
10-2
10-4
10-6
10-8
10-10
100
10-12
10-14
23
(Vo2+) or metal ions (m+) in the other direction. It is the accumulation of these defects in
metal oxide that may lead to the formation of filaments in the metal oxide film, and may
lead to conduction pathways [46]. A conductive path expansion due to drift of Vo2+
towards the top electrode interface from the applied negative voltage is called “virtual
cathode”. This “virtual cathode” will ultimately transform into a complete conducting
path when it reaches the bottom electrode and the device switches to the LRS [46,64]. To
switch the device back to the HRS an opposite bias voltage polarity is applied which
results in the dissolution of the formed filament. This electroforming process is
considered necessary for subsequent stable bipolar switching operations.
Figure 1.11: Gentle forming process – soft breakdown of oxide material.
-4 -3 -2 -1 0 1 2 3 4
Curr
ent (
I)
Voltage (V)
Gentle forming
100 nA500 nA
10 uA 250 uA
250 uA10-2
10-4
10-6
10-8
10-10
100
10-12
10-14
24
Mechanisms of resistive switching
The conductive filament model and interface model are two widely discussed
models to describe the resistive switching mechanisms for RRAM devices.
Conductive filament model
It has been proposed that oxygen ion defects, such as oxygen vacancies, in metal
oxides have a major role in resistive switching and have been considered to be much
more mobile than cations [65]. Following the electroforming process, bipolar switching
may take place through localized reduction/oxidation reactions. It is possible that these
redox reactions at the oxide and the metal electrode interface can cause the formation and
rupture of the conductive filaments. Understanding these conductive filaments began
formally when the filament model was first introduced early 1960’s [66]. Several
research groups have proposed that the mechanism behind resistive switching is the
formation and dissolution of conductive filaments in the oxide materials [67,68]. For
example, in 2009 Hangbing et al. investigated the resistive switching of Cu-oxide films
with several different top metal electrodes (Al, Pt, Ti ) [56]. It was found that the Al top
electrode showed good stability, endurance and a larger resistance ratio compared to the
other metals. After initial forming, LRS was achieved by applying a positive voltage on
the top electrode during the SET process. It was suggested that O2- ions were attracted to
the top electrode interface resulting in the formation of a thin AlOx layer. During the
RESET process the HRS was achieved by applying a negative voltage to the top
electrode where the O2- ions were repelled away from the top electrode interface which
25
are incorporated into the oxygen vacancies. Analysis using AES and TEM techniques
were performed to better understand the switching mechanism for these devices. The
TEM studies indicated the presence of a 5-7 nm AlOx layer at the Al-CuxO interface.
These results were consistent with AES results which showed enhanced signals of Al and
O near the Al-CuxO interface. Depending on the polarity of the applied voltage, the AlOx
layer acted as a supplier of oxygen or as an oxygen reservoir. This paper also reported
both unipolar behavior and proposed that this switching results from the formation and
dissolution of conductive filaments due to joule heating [59,69].
Measuring I-V characteristics using conductive atomic force microscopy (C-AFM) for
studying resistive switching in metal oxides has proven to be an effective method to
investigate switching in RRAM devices [70-72]. In 2009, Kim et al. investigated resistive
switching in NiO thin films by measuring its electrical properties by using C-AFM. The
data suggested that conductive filaments composed of metallic Ni clusters were generated
as a result of oxide reduction in NiO thin films from the external voltage biases [70]. A
general schematic for the forming, set and reset process associated with the conductive
filament model is shown in Figure 1.12.
Ffi
In
el
p
b
st
fa
igure 1.12: Filament form
nterface mo
The In
lectrochemic
olarity. The
arrier height
In 200
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abricated cel
Adapted fr
Forming, SEmation and ru
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resistive sw
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11, pp. 28, 2
leading to co
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RAM which
[75]. The st
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onductive
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73-76].
h had highly
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d by RF
26
bias
he
27
magnetron sputtering from a Ta target in an oxygen ambient. X-ray photoelectron
spectroscopy indicated that the material consisted of two phases of TaOx, mainly TaO2-β
in the bulk and Ta2O5-δ near the anode. During the RESET operation, when a positive
voltage was applied on the anode, the O2- ions migrated from the bulk TaOx layer
(oxygen reservoir) and accumulated at the anode leading to oxidation of TaO2-β to Ta2O5-
δ which enlarged the band gap of the oxide material and increased the interfacial barrier
height. This increase in the barrier height causes the device to be in the HRS. On the
other hand, in the SET operation the reduction of Ta2O5-δ occurs due to the migration of
O2- ions away from anode due to applied negative voltage. This process reduced the band
gap of the oxide material, and thereby decreasing the interfacial barrier height causing the
device to be in the LRS. The evidence of redox reaction happening at the interface was
confirmed by hard X-ray photoemission spectroscopy [75]. The spectra showed
increased signals from reduced components during the device’s transition from HRS to
LRS. Resistive switching following the same mechanism described above was previously
suggested by Muraoka et al. for FeO RRAM [74]. A general schematic for the SET and
RESET processes associated with the interface model is shown in Figure 1.13.
Fin
th
d
L
sw
Adapte
igure 1.13: Rnterface.
To un
he dependen
ependence o
LRS and HRS
witching me
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Resistive sw
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S was propo
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vol. 98, pp. 2
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2237, 2010
trode/oxide
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28
ng,
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29
1.6 Oxide RRAM Overview
The first practical application of RRAM was reported in 2002 by Zhuang et al.
with the fabrication of 64 - bit non-volatile memory based on perovskite oxide, PCMO
(Pr0.7Ca0.3MnO3), as the switching element using a 0.5 µm CMOS process [50]. At room
temperature the resistance ratio between the high resistance state to low resistance state
was larger than 1000. In 2003, Gould et al. evaluated the performance of aluminum top
and bottom electrodes on the resistive switching in comparision to gold top and bottom
electrodes for SiOx devices. This study suggested that electroforming took place with
these Al-SiOX-Al structures at the expense of more voltage cycles and the sample
stability with Al top electrodes was inferior in comparison to RRAM with gold as the top
electrodes, which could be due to comparatively lower melting point of Al [77].
Furthermore, the Poole – Frenkel effect was proposed to be the mechanism behind the
negative resistance observed in the SiOX films [77].
In 2004, Baek et al. demonstrated fabrication of RRAM based on a binary
transition metal oxide (TMO) switching material. The TMO was called OxRRAM and
was fabricated using 0.18 µm CMOS technology. The authors claimed that binary TMOs
are advantageous compared to perovskite oxides because of poor controllable of crystal
structure and stoichiometry for perovskite materials [51]. Their cells had a MIM structure
with polycrystalline TMO layer sandwiched between two noble metal electrodes. A
voltage pulse independent of voltage polarity was required to switch from a low
resistance to high resistance state and vice versa. The cell size dependency on the
30
resistance of the cells indicated that the SET resistance and RESET resistance decreased
only slightly with an increase in cell size which suggested that the switching mechanism
was due to the formation of filament current paths in the oxide film [51].
RRAM based on binary TMOs have attracted significant attention, especially for
resistance switching using TiO2. For example, Choi et al. demonstrated that RRAM
devices fabricated with TiO2 thin films with thicknesses of 20 nm to 57 nm could have
more than a few hundred continuous switching cycles with resistance ratios greater than
102. Conductive atomic force microscopy studies confirmed that the resistive switching
mechanism may be due to formation and elimination of conducting spots [52]. Shima et
al, have also shown that other binary TMOs have good performance for RRAM,
including Pt/NiO/Pt and Pt/CoO/Pt structures [53].
In 2007 Lin et al. studied resistive switching on doped films, in this case Mo –
doped SrZrO3 (SZO) [54]. SiO2/Si substrates were coated by RF magnetron sputtered
LaNiO3 that was the bottom electrode and 30 and 50nm SZO films were spincoated using
a solution containing strontium acetate, zirconium acetate and varying concentrations of
molybdenum acetate (0.1, 0.2, 0.3%) dissolved in acetic acid and acetylacetone. An Al
top electrode was evaporated and patterned using metal masks. I-V measurements
revealed good resistive switching for the SZO films doped with 0.1 and 0.2% Mo. The
resistance at the on and off state was stable over 104 s [54].
Xu et al., have studied the mechanism of resistive switching for TiN/ZnO/Pt
based RRAM devices, including dependence on cell area, operating temperatures, and
31
frequencies [56]. In this study ZnO was coated on a Pt electrode by reactive sputtering
and annealed at 450 °C for 30 min in O2/N2 ambient. XPS studies on the ZnO film
confirmed a high content of non-lattice oxygen ions in the material. Resistive switching
in the devices was observed without requiring the electroforming process. Based on the
change in resistance of the device on cell area and the temperature, it was concluded that
the current conduction in the device could be due to locally confined filaments and
electrons hopping through the oxygen vacancy defects [56].
Chen et al, have studied amorphous oxide materials for RRAM applications using
amorphous IGZO [60]. In this study the bottom electrode was 100 nm ITO coated on to a
glass substrate. The IGZO switching material was sputter deposited at room temperature
from a target containing 1:1:1 atomic ratio of In: Ga: Zn with a power of 50 W and
working pressure of 4 mTorr. Finally the ITO top electrode was sputter deposited through
a shadow mask with feature sizes of 100-800 µm on top of the IGZO film. For all these
measurements, the bottom electrode was grounded. The entire device was highly
transparent having 70 – 80% transmittance and XRD studies showed that the IGZO films
were amorphous. Stable bipolar resistive switching behavior was observed in the device
without an electroforming process. A compliance current of 10 mA was applied in the
SET process to prevent the device from breaking down. The resistance ratio between
HRS and LRS was found to be greater than 101. By plotting current density versus
voltage and determining the slopes at LRS and HRS, it was found that ohmic behavior
dominates the conduction in LRS whereas conduction in HRS was dominated by space-
charge-limited current. These studies strongly support the filament model [60].
32
1.7 References
[1] Y. Taur, T. H. Ning, “Fundamentals of Modern VLSI devices”, 1st Edition, Cambridge University Press (1998).
[2] R.C. Jaeger, “Introduction to Microelectronic fabrication”, 2nd Edition, Prentice Hall (2002).
[3] J. E. Lilienfeld, “Method and apparatus for controlling electric currents”, US patent, 1745175 (1930).
[4] J. E. Lilienfeld, “Amplifier for electric currents”, US patent, 1877140 (1932).
[5] J. E. Lilienfeld, “Device for controlling electric current”, US patent, 1900018 (1933).
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39
CHAPTER 2 - EXPERIMENTAL TECHNIQUES
This chapter covers various thin film deposition and patterning techniques
involved in fabrication of the TFT and RRAM devices. Additionally, material
characterization techniques for structural and compositional analysis of the amorphous
oxide material and TFT and RRAM device electrical characterization are discussed.
2.1Thin Film Deposition Techniques
Physical vapor deposition (PVD) refers to a group of vacuum-based techniques
used to deposit thin films by condensation of materials from the vapor phase onto solid
substrates. The two main PVD techniques used in this work are radio frequency (RF)
magnetron sputtering and thermal evaporation. In this section, the fundamental basis of
these techniques will be discussed, as well as important process variables that should be
considered when depositing films by these techniques.
2.1.1 Radio-Frequency Magnetron Sputtering
Sputter deposition is one of the most commonly used techniques for thin film
deposition. Sputtering occurs when a gas (usually argon) is ionized by an applied
potential resulting in a plasma. Due to electric and magnetic fields the ions are
accelerated towards the target, which consists of the desired material to be deposited.
When argon ions impinge on the target with high enough energy, a series of binary
collisions occur which results in the removal of atoms from the target material. These
ejected atoms can be transported from the target through the plasma, and are deposited on
40
the surface of the substrate. This process is allowed to take place for a given time until a
film of sputtered atoms is deposited on the substrate with the desired thickness. Sputter
deposition has several process advantages when compared to other deposition techniques,
including good film uniformity, high reproducibility, good control of film thickness, and
consistent film stoichiometry.
The RF plasma creates both positive and negative charged argon ions, although
positive charged ions are the dominant ionic species. The plasma is initiated by
application of a high RF voltage between a cathode and the anode with a low background
pressure of gas. The plasma in the system is sustained by ionization of the gaseous
species due to ion bombardment by secondary electrons emitted from the cathode.
Magnetron sources are employed to create a magnetic field that can be used to
trap secondary electrons close to the target. The electrons following helical paths due to
the magnetic field lines, and undergo more and more ionizing collisions near the target.
This cascading effect increases the ionization of gaseous species in the plasma, which
leads to an enhanced sputter rate from the target. A significant portion of the sputtered
atoms are not charged and are unaffected by the magnetic field from the magnetron.
RF sputtering was used in these studies due to the use of an insulating zinc tin
oxide target [1]. Since insulating targets tend to build up charge on the surface, an RF
potential can be used where an alternating negative/positive potential is applied to the
target electrode. During the first portion of the RF cycle ions are accelerated towards the
target surface when the target has a negative potential leading to sputtering of the target
m
th
th
2
F
in
sy
an
ro
pr
is
material, and
he target surf
he insulating
.1.
igure 2.1: Sc
An AJ
nvestigations
ystem is equ
nd one for 1
otation and t
ressure gaug
s obtained us
during the s
face when th
g target [2]. A
chematic of
JA Internatio
s and is loca
uipped with 4
inch targets
temperature
ges, and gate
sing a high v
second portio
he target has
A schematic
sputter depo
onal RF mag
ated at the M
4 sputter gun
s, both DC an
control, a lo
e valves. The
vacuum turbo
on of the RF
s a positive p
of a sputter
osition system
gnetron sputt
Microproducts
ns with autom
nd RF powe
oadlock, turb
e base pressu
o-molecular
F cycle electr
potential to m
r deposition s
m.
ter depositio
s Breakthrou
mated shutte
er supplies, a
bo-molecular
ure of the sy
r pump. The
rons are acce
minimize cha
system is sh
on system wa
ugh Institute
ers, three for
a substrate h
r pumps, dry
ystem is ~ 1x
instrument h
elerated tow
arge build-up
own in Figu
as used for t
e (MBI). The
r 3 inch targe
older with
y vacuum pu
x10-8 mbar w
has a gate va
41
wards
up on
ure
these
e
ets
umps,
which
alve
42
that separates the load lock (where the substrates are loaded) from the main chamber
(where the depositions are performed). The system can deposit on substrates as large as 6
inch diameter, using either 1 inch or 3 inch diameter targets. The system capabilities
include in-situ substrate heating, and can perform reactive and non-reactive sputter
deposition.
In this study both reactive and non-reactive sputter deposition were used to
deposit ZTO thin films for both TFT and RRAM applications where argon and oxygen
were used as the sputter gases.
2.1.2 Thermal Evaporation
Thermal evaporation is a physical vapor deposition technique in which a vapor of source
material is created by heating the material of interest to temperatures that increase the
materials vapor pressure. The evaporator used in this study was manufactured by Poloron
and is located in the OSU cleanroom in Owen Hall. Figure 2.2 shows a schematic of a
thermal evaporator.
The system is a high vacuum system and has a glass bell jar assembly. The system can
achieve a base pressure of 5 x 10-7 mbar using a diffusion pump. In these studies
aluminum (source material) is placed in a tungsten basket which is resistively heated by
passing a current of ~ 25 A through the basket. The aluminum slowly melts, subsequently
vaporizes, and finally condenses on to the substrate which is placed directly above the
source. In this study, the aluminum metal contacts for TFT source/drain and RRAM top
electrodes were deposited using thermal evaporation.
F
2
2
on
su
re
Adapted
igure 2.2: Sc
.2 Lithogr
.2.1 Shadow
Shado
n a substrate
ubstrate occu
elatively eas
d from Eric S
chematic of
raphy
w Mask Lith
ow mask lith
e. This metho
urs where th
y for pattern
Steven Sundh
thermal evap
hography
hography, als
od works by
here is an ope
ning source m
holm, MS th
porator syste
so called a st
y the selectiv
ening in the
materials, an
hesis, Oregon
em.
tencil metho
ve deposition
metal shado
nd skips man
n State Univ
od, was used
n of source m
ow masks. Th
ny of the pro
versity, 2010
d to pattern fi
material onto
his process i
ocessing step
43
0
films
o a
is
ps
44
used in photolithography. However, there are limitations in minimum feature size (>
1µm) which prevents its use in high precision applications [3]. In this study the ZTO
semiconductor layer and aluminum metal source-drain contacts for TFTs were patterned
by deposition of ZTO and aluminum through shadow mask sets. Our minimum feature
sizes using this technique was 100 µm gaps for the distance between source and drain for
TFTs.
2.2.2 Photolithography
Photolithography is a thin film patterning technique widely used in semiconductor
processing. This process involves transfer of patterns from a photo mask to a photoresist
(photosensitive organic polymer) under exposure to ultraviolet light [1]. The metal
electrodes and the ZTO for the RRAM devices were patterned using photolithography.
For a typical process positive S1818 photoresist was spin-coated on to the substrates. The
photoresist was soft baked at 90 ºC to remove solvents. A photomask is then placed
between the photoresist and an ultraviolet light source. The photomask prevents UV light
to interact with the photoresist in certain regions and allows the light to interact with the
photoresist in other regions. The interaction of UV light with the positive photoresist
drives photochemical reactions that cause the photoresist to become soluble in developer
solutions. The exposed regions of the positive photoresist were removed using Micro
Posit 351 developer. At this point the desired patterns as defined by photomask are left
behind. Pre-bake and post-bake of photoresist is done before and after development at 85
°C on a hot plate to avoid shrinking and cracking of photoresist. Typically, the films not
45
covered by photoresist are etched away. The photo resist is then removed from the
substrates by acetone, leaving behind the patterned film.
Another photolithography process that was used is lift-off where a material is directly
deposited on the substrate with patterned photoresist. The film is deposited on exposed
regions of the substrate and onto the resist. The undesired film was “lifted” away by
immersing the substrate in an acetone bath. Likewise, the film remains on the substrate
where there was no photoresist. The substrates were ultrasonicated to further remove any
unwanted film from the substrate. Lift-off was used as a replacement to shadow masks
and was useful when it is difficult to remove materials by etching due to poor etch
selectivity. One of the major disadvantages of lift-off patterning was that there can be
contamination at the interface between films since photoresist is deposited on to the
substrate prior to film deposition. In this work lift-off patterning of ZTO was performed
to pattern active switching material and aluminum was performed to pattern the metal top
electrodes for RRAM devices.
The photomasks used for the RRAM device fabrication process was designed using
AUTOCAD and the mask layout is shown in Figure 2.3.
Ffom
2
P
am
m
in
ch
re
an
re
al
[5
igure 2.3: Scor fabrication
mask and (C)
.3 Post-de
ost-depositio
morphous ox
material can b
n the electron
hannel mate
eduction in t
nnealed in a
esulting in d
lso helps imp
5]. The tube
chematic of n of RRAM ) Top electro
eposition a
on annealing
xide semicon
be increased
n carrier con
rials can be
the electron c
ir using a tub
iffusion of o
prove the se
furnace mak
the dark fielcross bar str
ode mask.
annealing
g is often use
nductor mat
d by annealin
ncentration in
reduced by a
carrier conce
be furnace. D
oxygen into t
emiconductor
kes use of re
ld and brightructures : (A
ed to modify
erials. For ex
ng in a reduc
n the materia
annealing in
entration [4]
During anne
the film that
r/dielectric i
esistive coils
t field photoA) Bottom el
y the electric
xample, the
cing atmosph
al, or conver
n an oxidizin
]. In this wor
ealing films r
t help passiv
interface by
s as heating e
omask sets uslectrode mas
cal propertie
conductivity
here resultin
rsely the con
ng atmospher
rk, the chann
react with O
ate defects.
reducing the
element and
sed in this stsk, (B) : Pad
s of the
y of the chan
ng in an incre
nductivity of
re resulting i
nel materials
O2 in the air,
The annealin
e number of
is equipped
46
tudy
nnel
ease
f the
in a
s are
ng
f traps
d with
47
a thermocouple for temperature monitoring. The samples were placed on a ceramic boat
during annealing in order to avoid contamination. A heating ramp rate of 10 °C/min and
dwell time of 1 hour was used for 300 and 600 °C post-deposition anneals.
2.4 Thin film characterization techniques
2.4.1 Spectroscopic Ellipsometry
Spectroscopic ellipsometry was used to determine the thickness of transparent films and
to obtain their optical constants by measuring the change in polarization of light from
either reflection or transmission as shown in Figure 2.4.
Figure 2.4: Ellipsometer for measurement of thickness and optical constants.
Thickness values ranging from sub- nanometer to several micrometers can be accurately
determined using this technique. The change in polarization is represented as Psi (ѱ) and
Polarizer Sample
Rotating analyzer
J.A. Woollam Co., Inc.
48
delta (∆) which depends on the optical properties and thickness of the films. Ѱ and ∆ are
related to the change in amplitude and phase of the polarized light by the following
equation,
= tan(ѱ) ∆
Where the terms Rp and Rs are the complex Fresnel reflection coefficients for the p- and
s- directions of the polarized light [6].
A material’s optical properties influence its interaction with light and can be modeled by
two values, the complex refractive index (ñ) and the complex dielectric function (ℰ). The
complex refractive index is represented by
ñ = −
where k is the extinction coefficient, which is related to absorption of light, and k=0 for
fully transparent materials and k>0 for absorbing materials. Likewise, n is the index that
characterizes the velocity of light, which is decreased as it enters a material with a higher
index of refraction. The complex dielectric function is represented by
ℰ = ℰ + ℰ
where ℰ1 is the real part of the dielectric function which is related to how strongly charge
in a material is polarized by the incident field, ℰ2 is proportional to amount of energy
absorbed from an applied field.
49
A regression analysis is performed using a model which allows one to compare extracted
information to the experimental data points. An initial estimate was used for the values of
thickness and optical constants for the film. Using these values a preliminary calculation
was performed using Fresnel’s equation. Depending on what is to be extracted, either the
film thickness or optical constants can be varied by an iterative process until the model
provided a good fit to the experimental data.
The mean squared error (MSE) is the measure of difference between the model and
experimental data [6]. The lower the value of MSE, the better the match between
experimental data and the model. The MSE function is
= 12 − ѱ −ѱѱ, + ∆ −∆∆, where a number of measured ѱ and ∆ pairs are represented by N, the total number of real
valued fit parameters are represented by M, and the standard deviations of ѱ and ∆ pairs
are represented by σexpѱ and σexp
∆ [6] .
There are two kinds of MSE minima called local minima and global minima which are
given by the regression model depending on how far the initial estimated value was from
the actual value, as shown in Figure 2.5 for an example thickness estimation. The global
minimum is the one which gives the correct thickness, which corresponds to the lowest
MSE value.
50
JA Woollam, Critical reviews of Optical Science and Technology, vol. CR72, pp. 3, 1999
Figure 2.5: Two kinds of MSE minima called local minima and global minima.
The extracted film thickness was correlated to the optical constants of a material which
will vary for different wavelengths. For transparent materials, the value of refractive
index is estimated using the Cauchy relationship,
( ) = + +
Where n represents the refractive index of the material, λ represents wavelength of light
and A, B and C are constants [3]. The Cauchy model was used in this work to determine
the thickness of various oxide films.
MSE
Thickness
Local Minima
Initialthickness estimation
Best Fit
51
2.4.2 Scanning Electron Microscopy
Scanning electron microscopy (SEM) is a technique that scans a beam of high energy
electrons (0.5 - 40 KeV) across a surface. This incident electron beam causes low energy
secondary electrons to be emitted from the surface of the sample and the secondary
electron current intensity is monitored with respect to the position of the scanned electron
beam [1]. Structure at the surface can be observed due to changes in the secondary
electron yield or secondary electron current intensity from the sample. For insulating
surfaces it is necessary to minimize charge build-up on the surface by the electron beam.
To minimize surface charging the corner of the sample was connected to a sample stub to
provide a conducting path from the surface and the surface can be coated with a thin
conductive film in addition to provide a means of charge dissipation. An example of an
SEM image is shown in Figure 2.6.
The magnitude of the secondary electron current strongly depends on the sample material
and sample topography, as there is more secondary electrons that escape from the
protruding edges that appear brighter in the image than recessed regions which appear
darker in the image. The image above was taken at a magnification of 25,000x and SEM
can provide magnifications of 500,000x [7].
F
2
X
b
p
sc
F
Figure 2.6: S
.4.3 X-ray D
X-ray diffrac
ased on the e
eriodic lattic
cattered X-ra
igure 2.7: Pr
d
X
SEM image o
Diffraction
ction (XRD)
elastic scatte
ce. Diffractio
ays from dif
rinciple of X
X-ray inciden
of platinum b
is a techniq
ering of X-ra
on occurs wh
fferent atomi
X-ray diffrac
A B
θ
nt beam
bottom elect
que used to id
ays by the el
hen there is
ic planes as s
tion.
Atoms in a c
B C
θ
trode of ZTO
dentify the a
lectron densi
constructive
shown in Fig
crystal
O RRAM cro
atomic struct
ity of individ
e interferenc
gure 2.7 [8].
Crystal Pla
oss bar struc
ture of a mat
dual atoms i
e between th
.
anes
52
cture.
terial
in a
he
53
The lattice spacing of the crystalline planes can be derived using Bragg’s relation
= 2
Where λ represents the X-ray wavelength, d represents spacing between the lattice
planes, θ represents the scattering angle, and n represents order of reflection where n is an
integer [8].
A wide variety of information can be obtained from XRD, including crystal structure,
lattice parameter, nanoparticle size, stress in the film, and an estimate on other properties
including disorders and imperfections in crystal structure.
The nanocrystallite size of a material can be determined using the Scherrer formula
< >=
where the nanocrystallite size (<L>), the peak width (β) and the scattering angle between
the incident beam and the angle normal to the reflecting plane (θ) are defined [8].
In this work, the XRD was used to determine the structure of the oxide thin films,
primarily the temperature for the amorphous to crystalline transition.
2.4.4 X-ray Photoelectron Spectroscopy
X-ray photoelectron spectroscopy (XPS) is a surface sensitive technique used to
determine atomic composition and oxidation state of material surfaces [9]. As shown in
Figure 2.8, this technique works by irradiating a sample surface with monochromatic X-
ra
T
ir
Fat
T
k
th
w
ty
T
=
fr
ays under va
This is illustr
rradiated wit
igure 2.8: Prtom (right).
The kinetic en
inetic energy
he photoelec
where the ene
ypically cons
The two most
= 1253.6 eV)
rom "L" shel
acuum. Thes
ated by the e
th photons th
rinciple of X
nergy of the
y (Ek) of the
ctron by the f
ergy of the X
stant for labo
t commonly
where Kα i
ll (principal
e X-rays cau
ejection of a
hat have high
XPS (left) an
se emitted p
emitted pho
following eq
X-rays (hν) a
oratory base
used x-ray s
s a X-ray em
quantum nu
use photo ele
a core-level e
her energy th
nd ejection of
hotoelectron
otoelectron i
quation,
= −and the work
ed instrumen
sources are A
mission line p
umber 2) to th
ectrons to be
electron from
han the bind
f a photoelec
ns is measur
is related to t
−
k function of
nts [10].
Al Kα (hν =
produced du
the "K" shell
e ejected from
m an atom w
ding energy o
ctron from in
red by the an
the binding e
f the spectrom
1486.6 eV)
ue to transitio
l (principal q
m the sampl
when it is
of the electro
nner shell of
nalyzer. The
energy (Eb)
meter (ϕ) is
and Mg Kα
ons of electr
quantum num
54
le.
on.
f an
of
α (hν
ron
mber
55
1). This transition defines the characteristic x-ray energy from the source. Insulating
samples need charge neutralization, which is performed by impinging low-energy
electrons or ions on the surface using a flood gun.
Figure. 2.9: XPS spectra for Zn 2p3/2 with binding energy and photoelectron intensity as X and Y axis.
As shown in the XPS plot in Figure 2.9, the elements present in the sample can be
identified from peaks at characteristic binding energies where in this example the Zn2p3/2
indicates the presence of Zn in the sample. For these studies the peaks were fit using
XPSPEAK4.1 software. The four main peak fitting parameters include peak position
(binding energy), peak area, peak full width at half maximum (FWHM) and %
Lorentzian – Gaussian (Line shapes). The background was approximated using a linear
approach.
0
10000
20000
30000
40000
50000
60000
1010 1015 1020 1025 1030 1035
Inte
nsity
(cou
nts)
Binding energy (eV)
Zn2p3/2 (hν = 1486.6 eV)
56
2.4.5 Rutherford Backscattering Spectrometry
RBS is an analytical technique widely used to determine the structure and composition of
materials. For RBS the sample is irradiated with a mono-energetic beam of high energy,
low mass helium ions, typically in the energy range of 1-3 MeV which gives the ions a
high penetration depth. The technique works by measuring the energy and number of
backscattered ions after colliding with atoms in the sample, the technique is
schematically shown in Figure 2.10. In this research the RBS technique was used to
determine the concentration of elements in the sample.
Figure 2.10: Principle of RBS.
For RBS, low mass incident ions are backscattered after undergoing elastic collisions
with the heavier atomic nuclei in the sample, resulting in a characteristic backscattered
energy. Furthermore, there is energy loss associated with the path length of the ions as
they travel through the sample. The RBS method is able to determine elemental
He+ Ion beam (2 MeV) 150°
He+
57
composition based on the energy of the backscattered ions, and the thickness and depth of
the film based on the width of the peak associated to a specific species. RBS is more
sensitive to detection of heavier atoms in comparison to lighter atoms because heavier
atoms tend to back scatter the incident ions more effectively.
An incident ion upon collision with surface atoms will backscatter at a different energy
than when it collides with a deeper atom in the lattice. When the collision of incident ion
with the deeper atoms takes place there is a greater loss of energy due to the incident ion
traveling through the matrix of atoms. This mechanism allows for depth profiling in RBS.
E1, the incident ion energy after collision is given by
= ( ± − ( ) )( + ) ∗
where m1 and m2 are the masses of incident ion and target atoms respectively, θ is the
scattering angle and E0 indicates incident ion energy [11]. The area under each individual
peak obtained from the analysis represents the total number of atoms of each element
present in a layer. By taking the ratio of the peak areas of two different elements, the
concentration ratio of the elements present in the film can be calculated with the
equation,
=
w
at
A
is
id
h
el
u
Fan
where C1, C2
tomic numbe
As shown in t
s plotted vers
dentified wit
igher energi
lastic collisio
sing the com
igure.2 11: Rnd backscatt
are concentr
er of two dif
the Figure 2
sus channel
th the charac
es in the spe
on theory [1
mmercial sof
RBS spectratering yield a
ration, A1, A
fferent eleme
.11, is an RB
number (ene
cteristic chan
ectrum due to
2]. In this w
ftware (SIMN
a of ZTO depas X and Y a
A2 are the are
ents present
BS spectrum
ergy). The el
nnel number
o the high va
work, the elem
NRA) [11].
posited on Siaxis.
ea under the
in the samp
m where the b
lements pres
s (energies).
alue for chan
mental comp
iO2/Si substr
peak, and Z
le.
backscatterin
sent in the sa
. Heavier ele
nnel energy
positions we
rates with ch
1, Z2 are the
ng yield (cou
ample can be
ements appe
due to binar
ere calculated
hannel numb
58
unts)
e
ar at
ry
d
ber
59
2.5 Electrical characterization using semiconductor parameter analyzer
All of the electrical characterization for TFTs and RRAMs were acquired using a probe
station with an Agilent 4155C precision semiconductor parameter analyzer. This
instrument measures data using four source monitor units (SMU's) (-40 V ≤ Vsource ≤
40V), (1 fA ≤ Isource ≤ 100mA). Three tungsten probes tips are used one each for the gate,
source and drain contacts for I-V measurements on TFTs and two tungsten tips are used
one for the bottom electrode and one for the top electrode for I-V measurements on
RRAMs.
60
2.6 References
[1] R.C. Jaeger, “Introduction to Microelectronic fabrication”, 2nd Edition, Prentice Hall(2002).
[2] D. M. Mattox, “Handbook of Physical Vapor Deposition (PVD) Processing”, 2nd Edition, Elsevier (2010).
[3] F. Ante , D. Kälblein , U. Zschieschang , T. W. Canzler , A. Werner , K. Takimiya , M. Ikeda , T. Sekitani , T. Someya , and H. Klauk , “Contact Doping and Ultrathin Gate Dielectrics for Nanoscale Organic Thin-Film Transistors”,small 7,1186 (2011).
[4] S.T. Meyers, J. T. Anderson, C. M. Hung, J. Thompson, J.F. Wager, and D. A. Keszler, “Aqueous Inorganic Inks for Low-Temperature Fabrication of ZnO TFTs”, Journal of American Chemical Society 130, 17603 (2008).
[5] H. Park, R. Choi, B.H. Lee, and H. Hwang, “Improved Hot Carrier Reliability Characteristics of Metal Oxide Semiconductor Field Effect Transistors with High-k Gate Dielectric by Using High Pressure Deuterium Post Metallization Annealing”, Japanese Journal of Applied Physics 46, L786(2007).
[6] J.A.Woollam, B.Johs, C.M.Herzinger, J.Hilfiker, R.Synowicki, and C.L.Bungay,“Overview of Variable Angle Spectroscopic Ellipsometry (VASE), Part I: Basic Theory and Typical Applications”, Critical Reviews of Optical Science and Technology CR72, 3 (1999).
[7] Scanning electron microscope - http://www.ece.umd.edu/class/enee416/GroupActivities/SEM%20vs%20FIB%20paper.pdf
[8] I.Chorkendorff, J.W. Niemantsverdriet, “Concepts of Modern Catalysis and
Catalysis”, 2nd Edition , Wiley-VCH (2007).
[9] D.A.Skoog, E.J.Holler, S.R.Crouch, “Principles of Instrumental analysis”, 6th edition, Thomson Brooks/Cole (2007).
[10] J.W.Robinson, E.M.S.Frame, G.M.Frame II, “Undergraduate Instrumental Analysis”, 6th edition, Marcel Dekker (2005).
[11] S.A.Chambers, M.H. Engelhard, V. Shutthanandan, Z. Zhu, T.C. Droubay, L. QiaO, P.V. Sushko, T. Feng, H.D. Lee, T. Gustafsson, A.B. Shah, J.-M. Zuo,Q.M. Ramasse, “Instability, intermixing and electronic structure at the epitaxial LaAlO3/SrTiO3(001) heterojunction”, Surface Science Reports 65, 317 (2010).
61
[12] J.L. Pau, L. Scheffler, M.J. Hernandez, M. Cervera, J. Piqueras, “Preparation of zinc tin oxide films by reactive magnetron sputtering of Zn on liquid Sn”, Thin Solid Films 518, 6752 (2010).
62
CHAPTER 3 – INVESTIGATION OF AMORPHOUS ZINC TIN OXIDE FILMS FOR THIN FILM TRANSISTOR APPLICATIONS
3.1 Introduction
Thin film transistors (TFTs) with transparent amorphous oxide semiconductors
(TAOS) represent a major advance in the field of thin film electronics.1 It has been
demonstrated that TAOS materials which contain heavy-metal cations with (n−1)d10 ns0
(n ≥4) electronic structure constitute a new class of high performance TFT channel
materials.1,2,3 It has been proposed that the high electron transport of TAOS materials is
due to the overlap of spherically symmetric empty metal s orbitals. These metal s orbitals
are unaffected by disorder due to the non-directionality of metal-metal s orbital overlap
and the nearly constant metal-metal distance in amorphous and crystalline materials.4
Consequently, these TAOS materials possess relatively high electron mobilities in spite
of being amorphous.3,4,5,6,7,8 Examples of TAOS materials include indium gallium zinc
oxide (IGZO),9 zinc tin oxide (ZTO),3 indium zinc oxide (IZO),10 and indium zinc tin
oxide (IZTO).11 Of these materials, ZTO does not contain indium or gallium and is
relatively inexpensive by comparison. ZTO has been used primarily as transparent
conductors12, as active material in sensors,13 and more commonly as a channel material
for TFTs.3,4,6,14 ZTO is a wide-bandgap (3.35 – 3.89 eV), n-type semiconductor material
which is transparent in the visible region of the electromagnetic spectrum4,6 and exists in
two crystalline phases: the cubic spinel Zn2SnO4 and the trigonal ilmenite ZnSnO3. 3,8 In
addition, ZTO has several good physical and chemical properties such as scratch
resistance, surface smoothness, and robustness to several chemical etchants.3,6
63
Several studies of ZTO TFTs have used RF magnetron sputter deposition, where
ceramic targets were used with a range of ZnO:SnO2 molar ratios.3,4 A combinatorial RF
sputtering technique has also been reported in which the ZTO channel layer was formed
by sputtering from two separate ZnO and SnO2 targets resulting in films with varying
Zn:Sn molar ratios.15 In these studies the performance of ZTO TFTs were investigated as
a function of channel stoichiometry3,4,15,16 and channel layer annealing temperature3,4.
McDowell et al. reported that TFTs with (ZnO)x(SnO2)1-x stoichiometry had the highest
values for electron mobilities at x = 0.25 and 0.80, although good performance was
observed over a broad range of stoichiometry.15
Chiang et al. reported that the mobility of the ZTO TFTs increased with
increasing annealing temperature and this could be due to the improved quality of the
channel – insulator interface with annealing. TFTs with channel layers annealed to 300
°C and 600 °C had channel mobilities ranging from 5 – 15 cm2/V.s and 20 – 50 cm2/V.s,
respectively. Hoffman recently demonstrated that the channel layer stoichiometry and
processing temperature strongly influenced a variety of device characteristics.4It wasshown that the turn on voltage (Von), on-off ratio, incremental mobility, and device
hysteresis were all sensitive to processing and composition. The highest values of
incremental mobility were ~ 25 to 30 cm2 V-1 s-1 which was obtained for ZTO films
formed with intermediate compositions of Zn and Sn and post- annealed to 400 – 600°C.
Several research groups have also studied the effect of oxygen partial pressure17
and doping 18 on ZTO films. Both optical and electronic properties of amorphous ZTO
films can change significantly due to differences in oxygen partial pressure during
64
deposition. Oxygen partial pressures of 2 – 7 Pa resulted in high carrier concentrations
and mobilities.17 Satoh et al. have also shown that conductivity and free carrier
concentration can be modified by incorporating Al in ZTO films.18 Other methods for depositing ZTO films have been demonstrated where several
groups have used solution-based methods including spin coating, 19, 20 ink jet printing21, 22
and dip coating.23 Although these methods are simple and may eventually lead to a low
cost path to manufacture TFTs, we have used sputter deposition since it allows excellent
control over the electrical and optical properties by varying pressure, power, and oxygen
partial pressure.
In this paper we report the physical and electrical characterization of ZTO films
formed by RF sputter deposition. These films were deposited at room temperature with
constant RF power where the argon/oxygen ratios were varied during deposition. Film
and device characteristics were evaluated for ZTO films following post-deposition
annealing.
3.2 Experimental details
All TFTs were prepared using test structures consisting of heavily doped p-type
silicon coupons that served as the gate and had a 140 nm thick thermal oxide that served
as the gate dielectric. Prior to ZTO deposition the substrates were sonicated in acetone,
methanol and iso-propyl alcohol for 5 min each and then rinsed with DI water. The
substrates were then blow dried using nitrogen and heated on a hot plate for 5 min at 150
°C to remove moisture prior to deposition.
65
The ZTO channel layer was deposited via RF magnetron sputter deposition using
a 3 inch ZTO target. The ZTO target was purchased from AJA International Inc. with a
ZnO:SnO2 ratio of 2:1 which represents a stoichiometry of Zn2SnO4. The ZTO films were
deposited with a constant 20 SCCM flow rate, but with different O2:Ar ratios
corresponding to 0:20, 1:19, 2:18 SCCM. Throughout the deposition process, the power
was maintained at 100W with a chamber pressure of 4 mTorr. Substrate rotation was
used to improve film uniformity and deposition times were adjusted to obtain the desired
thickness of 50 nm, which was confirmed by spectroscopic ellipsometry. The channel
layer was patterned using a shadow mask during the deposition process. The ZTO films
were post-annealed in air with a ramp rate of 10 °C/min and 1 hour dwell time at either
300 °C or 600 °C.
TFTs were fabricated using a staggered bottom gate structure, where source and
drain electrodes were patterned by depositing ~500 nm of Al via thermal evaporation
through a shadow mask. The width/length (W/L) ratios of the fabricated devices were
1000 µm/200 µm and 1000 µm/100 µm. TFT transfer characteristics were measured
using a 4155C Agilent semiconductor parameter analyzer and a probe station.
Scanning electron microscope (SEM) imaging was performed on the as-deposited
samples using a FEI Helios Nanolab dual-beam focused ion beam/scanning electron
microscope (FIB/SEM). The ZTO samples studied in this investigation had an insulating
surface that caused charge build-up on the surface by the electron beam. To minimize
66
charge build-up the corner of the sample surface was connected to the sample stub using
carbon tape to provide a conduction pathway.
X-ray diffraction (XRD) measurements were performed to identify the structure
of the films obtained under different sputter and annealing conditions. The
measurements were performed using a Rigaku Rapid image-plate system with a rotating
anode Cr Kα radiation source that was collimated using a 0.3 mm pin-hole and a 15°
incident angle.
Surface and bulk compositional analysis was performed using X-ray
photoelectron spectroscopy (XPS) and Rutherford backscattering spectrometry (RBS),
respectively. The PHI 5000 VersaProbe XPS system used monochromatic Al Kα
radiation with a photon energy of 1486.6 eV. The analyzer was operated with a pass
energy of 23.5 eV and all spectra were taken at normal emission. An electron flood gun
was used for charge compensation for the insulting samples. XPS sputter depth profiles
were obtained using an Ar ion beam (2 keV) on a 1x1 mm2 area for ZTO films deposited
with 1 SCCM of oxygen at room temperature. For the XPS experiments from annealed
samples, the samples were placed into the loadlock directly after annealing once they
were ≤ 100 °C to minimize ambient surface contamination. RBS data was acquired using
a high energy helium ion beam (2 MeV) system that has been described previously.24 For
these studies we used an incident angle of 90° and a backscatter angle of 150°. The RBS
data was analyzed using SIMNRA. 24
67
3.3 Results and discussion
In Figure 3.1 we show XRD data obtained from as deposited ZTO films and after
annealing to 600 °C and 650 °C. There is a lack of strong diffraction peaks in the spectra
for both the as deposited and 600 °C films, indicating that the films are amorphous even
after annealing up to these temperatures.3 These results are similar to those published
previously for ZTO films where broad peaks at 34° and 56° are representative of
amorphous ZTO.1,16,20,25,26,27,28 Also shown in Figure 3.1 are the location of major
diffraction peaks for both ilmenite ZnSnO3 and spinel Zn2SnO4.3
After annealing the ZTO films to 650° C we found that the ZTO crystallized,
which is evident from the sharp XRD peaks shown in Figure 3.1. These results are
similar to prior literature data,3,25,29,30 where the broad peak at 52.33° corresponds to the
(311) Zn2SnO4 diffraction plane while the peak at 38.5° corresponds to the (200) SnO2
diffraction plane. We found that the spinel Zn2SnO4 was the dominant crystalline phase
with trace amounts of SnO2 after annealing at 650 °C which matches the results from
several other authors and is expected based on the stoichiometry of the films.31,32,33
F6
a
n
th
ab
igure 3.1: X50° C.
In Fig
magnificatio
anoscale stru
hat the as dep
bove.
XRD spectra
gure 3.2 we s
on of 65,000
ucture that w
posited film
for ZTO thin
show an SEM
0x. The film
would be con
s are amorph
n films as-de
M image of t
is found to b
nsistent with
hous, in agre
eposited and
the as-depos
be very smo
h crystallizati
eement with
d annealed to
sited ZTO fil
ooth and doe
ion. These r
the XRD re
o 600° C and
lm obtained
es not show
esults sugge
esults discuss
68
d
with
est
sed
F
Z
th
fo
sp
sp
th
b
w
p
ph
igure 3.2: SE
XPS w
ZTO films du
he main low
or any charg
pectra had p
pectra were
hese spectra
ackground r
width at half
eak position
hotoelectron
EM image o
was used to i
ue to change
energy carb
ge induced sh
eaks due to z
obtained for
peak fitting
removed and
maximum (F
ns were adjus
n peaks were
of as-deposite
investigate t
s in depositi
bon 1s peak (
hifts since th
zinc, tin, oxy
r the Zn 2p3/2
was perform
d were then f
FWHM), %L
sted during t
e obtained fo
ed ZTO film
the changes i
ion and post-
(aliphatic ca
he ZTO films
ygen, and ca
2, Sn 3d5/2, O
med using XP
fitted with a
Lorentzian –
the peak fitti
or the films a
m taken at 65
in chemical
-processing
arbon) was se
s are insulati
arbon for the
O 1s, and C 1
PSPEAK4.1
linear backg
– Gaussian (%
ing procedur
annealed to 6
5,000x magn
state and co
conditions. T
et to 284.6 e
ing.34 The X
e ZTO films.
1s photoelec
1.34 The peak
ground and t
%L-G) and b
re. The shar
600° C and t
nification.
omposition o
The position
eV to compen
XPS survey
. High resolu
ctron peaks.
ks had their
the area, full
binding ener
rpest
these data w
69
of the
n of
nsate
ution
For
l
rgy
were
u
te
ob
eV
3
an
d
st
Fan
T
ti
sed to define
emperature f
btained the f
V and 40% a
d5/2 spectra f
nd 600 °C. A
etermined to
tates of Zn2+
igure 3.3: Xnnealed to 3
These binding
in. 11,19, 35 Fur
e the initial f
films were u
following FW
and 45%, res
for as deposi
After charge
o be 1021.49
+ and Sn4+, re
XPS of Zn 200 °C and 6
g energies ar
rthermore, w
fitting param
sed for the f
WHM and %
spectively. I
ited films wi
compensati
9 eV and 486
espectively.
2p3/2 from Z00 °C.
re close to pu
we found tha
meters, excep
fitting param
%L-G values
In Figure 3.3
ith 1 SCCM
on the bindi
6.22 eV, whi
ZTO films d
ublished val
at the Zn2p3/2
pt for the C 1
meters. Based
s for Zn2p3/2
3 and 3.4 we
M of oxygen a
ing energy fo
ich correspo
deposited w
lues for oxid
2, and Sn 3d
1s peak whe
d on this ana
, and Sn 3d5
e show the Z
and films an
or Zn2p3/2, a
onds to Zn an
with 1 SCCM
de films cont
d5/2 binding e
re the room
alysis we
/2 1.65 and
Zn2p3/2, and
nealed to 30
and Sn 3d5/2 w
nd Sn oxidat
M of oxygen
taining zinc
energies wer
70
1.38
Sn
00° C
were
tion
n and
and
re
n
an
Fan
fi
fi
h
w
1
T
n
ox
early identic
nnealing con
igure 3.4: Xnnealed to 3
In Fig
ilms prepare
itting parame
igh energy s
which were d
s spectra we
The lower en
eighboring Z
xygen from
cal for films
nditions.
XPS of Sn 3d00 °C and 6
gure 3.5 we s
d with an ox
eters for the
shoulder did
determined to
e found that i
ergy peak (I
Zn and Sn io
a hydroxide
deposited w
d5/2 from ZTO00 °C.
show XPS O
xygen flow r
O 1s peak w
not significa
o be 1.42 eV
it was neces
I) at 530.02 e
ons. The high
e species pre
with 0, 1 and
O films depo
O 1s spectra o
rate of 1 SCC
we again use
antly influen
V and 10%, r
sary to use t
eV correspon
her energy p
sent on the s
2 SCCM of
osited with 1
of the as-dep
CM during d
ed the 600° C
nce the value
respectively.
three peaks t
nds to lattice
peak (II) at 5
surface and i
f oxygen and
SCCM of o
posited and a
deposition. T
C annealed fi
es for FWHM
Using these
to adequately
e oxygen ion
31.50 eV ca
in the films.
d irrespective
oxygen and
annealed ZT
To obtain pea
films since th
M and %L-G
e values for t
y fit the spec
ns, which ha
an be assigne
The intensit
71
e of
TO
ak
he
G
the O
ctra.
ave
ed to
ty of
th
(I
fi
sl
su
O
th
fi
re
Fan
his peak was
III) at 532.52
ilm.11,19,36 Th
lightly highe
ubstrate peak
OH- and H2O
hat these hig
ilms and this
ecombination
igure 3.5: Xnnealed to 3
s significantl
2 eV can be
he FWHM o
er FWHM (1
k. This is to
O for the film
gher energy p
s is likely du
n of hydroxy
XPS of O 1s f00 °C and 6
ly reduced af
assigned to
of the hydrox
1.47 eV) com
account for
ms, and this li
peak intensit
ue to desorpti
yl species or
from ZTO fi00 °C.
fter annealin
oxygen from
xyl (II) and w
mpared to the
any inhomo
ikewise imp
ties were sig
ion of water
r via molecu
ilms deposite
ng the films.
m water pres
water (III) p
e O 1s peak
ogeneity in th
roved the qu
gnificantly re
r from the su
ular desorptio
ed with 1 SC
The highest
sent on the su
peaks were a
correspondi
he chemical
uality of the
educed after
urface throug
on.37
CCM of oxy
t energy pea
urface and in
llowed to ha
ing to the
bonding for
fit. We foun
annealing th
gh the
gen and
72
ak
n the
ave a
r the
nd
he
fi
h
th
p
an
fo
is
2
su
Fan
In Fig
ilms prepare
ad three fair
he films. To
eak for the u
nd 40%, resp
or the three p
s assigned to
86.21 eV (II
urface, respe
igure 3.6: Xnnealed to 3
gure 3.6 we s
d under an o
rly well sepa
obtain peak
unannealed f
pectively. Pe
peaks, with t
o aliphatic ca
I) and 288.57
ectively.38
XPS of C 1s f00 °C and 6
show C 1s X
oxygen flow
arated peaks,
fitting param
films and the
eak fitting of
the lower en
arbon (C-C o
7 eV (III) ca
from ZTO fi00 °C.
XPS spectra o
rate of 1 SC
where all th
meters for th
e FWHM an
f C 1s spectr
nergy peak ha
or C-H bond
an be assigne
lms as-depo
of the as-dep
CCM during
hree reduced
he C 1s peak
nd %L-G wer
ra allowed u
aving an ene
ds), and the tw
ed to C-O an
osited with 1
posited and a
deposition.
d in intensity
k we used the
re determine
us to extract b
ergy of 284.6
wo higher en
nd C=O grou
SCCM of o
annealed ZT
The C 1s sp
y after annea
e low energy
ed to be 1.32
binding ener
6 eV (I), wh
nergy peaks
ups on the
oxygen and
73
TO
pectra
aling
y
2 eV
rgies
hich
at
74
Table 3.1 gives the binding energy values for Zn 2p3/2, Sn 3d5/2, O 1s, and C 1s
for samples deposited with different O2 flow rate and post-annealing conditions. The
binding energy shifts for each condition varied < 0.1 eV. The atomic compositions of Zn,
Sn, O in the ZTO films were estimated using the area under the individual peaks and the
relative sensitivity factors, and the atomic concentrations are given in Table 3.2. The
Zn/Sn and O/(Sn +Zn) ratios were determined to be 1.61 and 1.47 for the as-deposited
ZTO film prepared with 1 SCCM of oxygen during deposition, which is close to the
target stoichiometry which are 2 and 1.33 respectively. On average, the atomic
concentrations of Zn, Sn and O were consistent regardless of changes in oxygen flow rate
during deposition.
Table 3.1: XPS table of binding energies for Zn 2p3/2, Sn 3d5/2, O 1s and C 1s for ZTO films with different deposition conditions and annealing temperatures.
CONDITIONS O1s (I) O1s (II) O1s (III) Zn 2p3/2 Sn 3d5/2 C1s (I) C1s (II) C1s (III)
0 sccm - RT 529.96 531.42 532.37 1021.53 486.22 284.60 286.27 288.590 sccm - 300 °C 530.12 531.61 532.59 1021.59 486.40 284.60 286.17 288.490 sccm - 600 °C 529.98 531.50 532.44 1021.46 486.16 284.60 286.16 288.751 sccm - RT 530.02 531.50 532.52 1021.49 486.22 284.60 286.21 288.571 sccm - 300 °C 530.00 531.40 532.46 1021.44 486.23 284.60 286.19 288.941 sccm - 600 °C 530.03 531.46 532.49 1021.46 486.23 284.60 286.20 288.752 sccm - RT 529.90 531.30 532.31 1021.30 486.12 284.60 286.15 288.662 sccm - 300 °C 529.95 531.47 532.39 1021.44 486.15 284.60 286.29 288.252 sccm - 600 °C 529.92 531.37 532.37 1021.31 486.13 284.60 --- ---
BINDING ENERGY (eV)
T
to
su
T
W
th
F
Table 3.2: XP
In Fig
o just before
urface with a
The surface e
We found tha
hroughout th
igure 3.7: X
ZTO AS
VARIO
PS table of a
gure 3.7 we s
the ZTO an
a Zn/Sn ratio
enrichment o
at annealing
he film comp
XPS sputter d
S-DEPOSITED FILMS W
US OXYGEN PERCENTA0 sccm1 sccm2 sccm
atomic ratios
show results
nd SiO2 inter
o of ~1.13, w
of ZTO films
films up to 6
pared to the a
depth profile
WITH EXPERIMENTAAGES
1.681.931.79
for as-depo
from an XP
rface. We fou
while the bul
s has been ob
600 °C resul
as deposited
of ZTO film
AL Zn/Sn EXPERIMENT
1.1.1.
osited ZTO th
PS sputter de
und that the
lk of the film
bserved by o
lted in minor
d film as show
m deposited
TAL O/(Zn+Sn) IDEAL Z
.59 2.
.64
.54
hin films.
epth profile f
ZTO film is
m has a Zn/S
other researc
r changes in
wn in Figure
with 1 SCCM
Zn/Sn IDEAL O/(Zn+
.00 1.33
from the surf
s tin rich at t
Sn ratio of ~1
chers as well
n composition
e 3.7.
M of oxygen
+Sn) IDEAL O/(Zn+Sn)EXPERIMENT
75
face
the
1.89.
l.39
n
n.
) CORRESPONDING TOTAL VALUE of Zn/Sn
1.251.321.28
Z
S
h
fo
la
th
sc
ca
ca
w
F
In Fig
ZTO films wi
n are found
igher atomic
ound in lowe
ayer and the
In the
hickness and
cattering ang
alibration va
alculate the b
were varied to
igure 3.8: R
gure 3.8 we s
ith 0 SCCM
at higher ch
c mass. Like
er energy reg
Si substrate
SIMNRA s
d compositio
gle, atomic n
alues for dete
backscattere
o match the
RBS spectra f
show experim
of oxygen.
hannel numbe
wise, Si and
gion of the s
.40
imulation, th
n. Experime
number of in
ectors were e
ed spectra. T
simulated an
from an as-d
mental and s
The features
ers, which c
d O have low
pectrum. Th
he sample is
ental parame
ncident ion, s
entered into
The values of
nd experime
deposited ZT
simulated RB
s are labeled
orrespond to
wer atomic m
he signals for
s divided into
eters such as
solid angle o
the SIMNR
f film compo
ental spectra.
TO film with
BS spectra f
d in the spect
o higher ener
mass, so their
r Si are due
o several lay
incident ion
of detection a
RA simulatio
osition and a
.24
h 0 SCCM of
for as-deposi
tra, where Zn
rgy, due to t
r signals are
to both the S
yers with var
n energy, ion
and energy
on software t
areal densitie
f oxygen.
76
ited
n and
their
SiO2
rying
n
to
es
77
The RBS data indicates that the as-deposited film had atomic percentages of Zn
and Sn close to that of target composition and the XPS data. Table 3.3 shows the ratios of
Zn/Sn for the as-deposited ZTO films with 0 SCCM, 1 SCCM and 2 SCCM of oxygen
during deposition. No significant change in film stoichiometry was observed between
ZTO films but all the values differ slightly from the nominal target composition.
Table 3.3: Fitting parameters for RBS spectra obtained using simulation
In Figures 3.9, 3.10, and 3.11 we show the transfer characteristics for ZTO TFTs
that have not been annealed or have been annealed to 300° C or 600° C, respectively. For
these measurements the drain current (ID) and gate current (IG) were measured while the
gate to source voltage (VGS) was scanned from -20 V to 20 V in a double sweep mode
where the drain to source voltage (VDS) was held constant at 1 V. In this work we use the
definition of the turn-on-voltage (Von) which is the gate voltage at which there is sharp
increase in ID in the log ID -VGS transfer curves. Other parameters of interest include the
ID on to off ratio and channel mobility.1,41
0 sccm 3.00 X 1017 1.601 sccm 2.90 X 1017 1.612 sccm 2.90 X 1017 1.80
ZTO AS - DEPOSITED FILMS WITH VARIOUS OXYGEN PERCENTAGE FILM THICKNESS (at/cm2) Zn/Sn Zn/Sn
FS
FS
igure 3.9: TFCCM, 1 SCC
igure 3.10: TCCM, 1 SCC
FT transfer cCM, 2 SCCM
TFT transferCM, 2 SCCM
characteristiM of oxygen
r characteristM of oxygen
cs with as-dn.
tics with ZTn and anneal
deposited ZT
TO channel laled to 300 °C
TO channel la
ayer depositC.
ayer with 0
ted with 0
78
FS
T
W
le
tr
ox
O
igure 3.11: CCM, 1 SCC
The value of
Where GCH is
ength ratio o
As sho
ransistor cha
xygen were
OFF characte
TFT transfCM, 2 SCCM
average chan
s the value o
f the channe
own in Figu
aracteristics.
conducting
eristics, wher
fer characteM of oxygen
nnel mobilit
μ (of channel co
el, and Cins th
ure 3.9 the as
We found th
which is evi
reas when ox
eristics withn and anneal
ty was extrac
) = onductance w
he capacitan
s-deposited Z
hat the ZTO
ident from th
xygen was a
h ZTO chanled to 600 °C
cted using,
( )( −with respect
nce of the SiO
ZTO TFTs d
films depos
he high drain
added during
nnel layer dC.
) to VGS, W/L
O2 insulator.
do not exhibi
sited with 0 S
n current wit
g deposition
deposited w
L is the width
.40
it typical
SCCM of
th no ON an
the films be
79
with 0
h to
nd
ecame
80
insulating which is evident from the low drain currents. After annealing the ZTO to 300°
and 600° C we have found that the TFTs have negative turn-on voltages of approximately
-15 V to -5 V depending on chamber oxygen flow rates and annealing conditions. This
suggests that carrier concentration can be controlled by the O2 partial pressure during
deposition where higher carrier concentration shifts the turn-on voltage more negative.
Similar results were also observed previously by Hong et al .6 Another important
parameter that determines the switching quality of TFTs is the drain current on-to-off
ratio.8 Both Figure 3.10 and 3.11 show that the drain current on-to-off ratio can be greater
than 106 for a broad range of processing conditions and their values increased with
increasing anneal temperature. Channel mobility is another parameter that quantifies the
performance of the semiconductor channel.3 The average mobility (µavg) corresponds to
the average mobility of all carriers in the semiconductor channel. We have found that the
average mobility for a given oxygen flow rate increases with increase in temperature due
to a variety of reasons, one of which is the modification of the semiconductor-insulator
interface and the change in the nature of traps in the semiconductor.3,4 It was observed
that the average channel mobility changed with respect to oxygen flow rate with two
different trends. As shown in Figure 3.12 the ZTO samples annealed at 300 °C had a
significant decrease in µavg with increasing oxygen flow rate, while at 600 °C there was a
slight increase in µavg with increasing oxygen flow rate. These results indicate that the
ZTO films electrical properties are still strongly influenced by the film deposition
conditions for lower annealing temperatures, while the initial deposition conditions are
le
cm
an
Fp
hy
h
d
o
b
ess importan
m2/V·s) wer
nnealed to 6
igure 3.12: ost-annealin
Furthe
ysteresis bet
ave the abili
We ha
evice charac
f the ZTO fi
oth the depo
nt for higher
e obtained fo
00 °C.
Change in ang conditions
ermore, after
tween the sw
ity to stay in
ave found th
cteristics and
ilms. In gene
osition condi
annealing te
for ZTO TFT
average chans.
r a 600 °C an
weep up and
n equilibrium
hat annealing
d the relative
eral the elect
itions and an
emperatures.
Ts sputter dep
nnel mobilit
nneal we fou
down traces
m with the sw
g the ZTO fil
e concentrati
trical propert
nnealing tem
.42 High aver
posited with
ty as a funct
und that the
s, suggesting
weep rate of
lms resulted
ion of oxyge
ties of the fi
mperature, wh
rage channel
h 2 SCCM of
tion of oxyg
TFTs do no
g that if traps
applied volt
d in significa
en impurities
ilms can be m
here ZTO ele
l mobilities (
f oxygen and
gen flow rate
t exhibit
s are present
age.43
ant changes i
s on the surfa
modifying by
ectrical
81
(17
d
e and
t they
in the
face
y
82
characteristics can be tailored to behave as insulating, semiconducting, or conducting.
High average mobilities have been obtained for ZTO TFTs suggesting that TAOS
materials can be considered as a replacement for amorphous silicon TFTs for next
generation displays.
3.4 Conclusion
We have analyzed the role of both deposition and post annealing conditions on
ZTO film structure, composition, surface contamination, and TFT device performance.
We found that the ZTO thin films were amorphous even after annealing to 600 °C and
that the as-deposited ZTO film composition was close to the target stoichiometry of
Zn2SnO4. XPS indicated that the ZTO films had significant surface contamination, and
that these could be reduced by higher temperature anneals. ZTO TFTs with a high
channel mobility of ~17 cm2/V·s were fabricated for sputter deposited ZTO films after
annealing to 600 °C. Device results also indicate that the electronic properties are closely
linked to deposition conditions, surface contamination, and post-processing conditions.
83
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86
CHAPTER 4 – BIPOLAR RESISTIVE SWITCHING IN SPUTTER DEPOSITED AMORPHOUS ZINC TIN OXIDE DEVICES
4.1 Introduction
Resistance random access memory (RRAM) devices have the potential to replace
silicon-based flash memory or other advanced memory architectures due to the simple
device structure, high areal density, low power consumption, and compatibility with
complementary metal oxide semiconductor (CMOS) technologies [1]-[4]. It has recently
been argued that the non-volatile switching of RRAM devices can be correlated with
memristors where the resistance of the device can depend on the charge that passes
through it [4, 5]. These devices are fabricated using metal-insulator-metal (MIM)
structures where the devices can be switched between a high resistance state (HRS) and a
low resistance state (LRS) depending on polarity of applied voltages and the initial state
of the device. The observation of large negative resistance in oxide thin films, including
SiOx, Ta2O5, ZrO2, Al2O3 and TiO2, was first reported in 1962 [6]. However it was not
until much later before oxide materials were incorporated into switching devices with
controlled structure and electronic properties [1, 7]. In the literature many oxide based
materials have been shown to have resistive switching including NiO[2], TiO2[8],
CoO[3], SrZrO3[9], Al2O3[10], ZnO[11], ZrO2[12], HfOX[13], CuXO [14] [15], TaOx
[16], indium gallium zinc oxide (IGZO) [17] and gallium zinc oxide (GZO) [18].
Recently, transparent amorphous oxide semiconductor (TAOS) materials have
been used as resistive switching materials [19]. These materials contain heavy-metal
cations with (n−1) d10 ns0 (n ≥4) electronic structure. Unlike hybridized covalent
87
semiconductors, the electron conduction band is made up of large, spherically
symmetrical metal s orbitals, and their overlap is unaffected due to structural disorder
[20]. These TAOS materials provide several benefits for RRAM devices including low
temperature processing, high uniformity, and atomic-scale surface smoothness. One
promising TAOS material is zinc tin oxide (ZTO), a wide band gap n-type semiconductor
material which is relatively inexpensive when compared to IGZO which contains
increasingly expensive gallium and indium based materials. ZTO has several good
physical and chemical properties such as high scratch resistance and robustness to several
chemical etchants [21], [22].
Recently sputter deposited IGZO was evaluated for RRAM applications where a
transparent indium tin oxide (ITO)/IGZO/ITO structure was fabricated at room
temperature [17]. Stable bipolar resistance switching was observed without the need for
an initial electroforming step and the switching ratio between the high resistance state
(HRS) and low resistance state (LRS) (RHRS/RLRS) was ~30 after ~100 switching cycles.
The conduction mechanisms for the LRS and HRS were dominated by Ohmic and space
charge limited current (SCLC), respectively. The authors proposed that the observed
switching was consistent with a conductive filament model. An all solution-processed
transparent RRAM with an ITO/GZO/ITO structure has recently been demonstrated [18].
For these devices good switching endurance was observed, however a fairly low
RHRS/RLRS ratio of ~15 was obtained. These devices also did not require a forming step
prior to switching and it was proposed that the main conduction mechanism during the set
process is due to trap-controlled space-charge limited current, where the resistance
88
change is due to modulation of the potential barrier height due to Fowler-Nordheim
tunneling.
In this study, we fabricate amorphous ZTO based RRAM and demonstrate bipolar
resistance. We have found that the switching properties of ZTO were closely related to
the electrical test conditions and the electrode materials. The physical and electrical
characterization of the ZTO materials and devices, along with the switching mechanisms
will be discussed.
4.2 Experiments
Memristor cross bar structures were fabricated using lift-off processing. The
bottom electrode was 50 nm thick platinum with a 25 nm thick titanium adhesion layer.
Both layers were deposited using an e-beam evaporation without removing from vacuum
between depositions, and the substrate was a silicon wafer with an insulating SiO2 layer.
Just prior to the ZTO deposition the substrates were rinsed with acetone, isopropyl
alcohol, DI water and then blow dried using nitrogen, and heated for 5 min at 100 °C on a
hot plate in order to remove absorbed moisture. A 50 nm ZTO layer was deposited via
RF magnetron sputter deposition using a 3-inch ZTO target with a ZnO : SnO2 ratio of
2:1 (Zn2SnO4). The ceramic target was purchased from AJA International Inc. The ZTO
films were deposited with 100 W power, a 20 SCCM flow rate with a 2:18 O2:Ar ratio,
and ~ 4 mTorr chamber pressure. The deposition time was adjusted to obtain a ~ 50 nm
thick film which was measured using spectroscopic ellipsometry. Substrate rotation was
used during deposition to improve film uniformity. No post-deposition annealing of the
ZTO films was performed. The aluminum top electrodes were also patterned by lift-off
89
where ~200 nm Al was deposited by thermal evaporation onto the ZTO film. The cross
bar structures had widths of 10, 20, 50 and 100 µm leading to cross bar cell areas of 100,
400, 2500 and 10000 µm2.
X-ray diffraction (XRD) measurements were performed using a Rigaku Rapid
image-plate system with Cr Kα radiation, a 0.3 mm pin-hole collimator, and a 15°
incident angle provided by a rotating anode. XRD was used to confirm that the as
deposited films were amorphous. RBS measurements were performed on the as-deposited
films using a high-energy helium ion beam (2 MeV), an incident angle of 90°, and a
backscatter angle of 150°. The data was analyzed using SIMNRA. RBS was used to
confirm the stoichiometric composition of the ZTO films. Sputter depth profiles were
obtained from the Al/ZTO/Pt stack using an ION-TOF IV Time-of-Flight Secondary Ion
Mass Spectrometry (TOF-SIMS) using 1 keV Cs+ ions for the sputter depth profile and
25 keV Bi3+ ions for analysis. The positive secondary ions were analyzed for these
studies. The electrical switching characteristics were measured using a 4155C Agilent
semiconductor parameter analyzer and a probe station. All electrical measurements were
performed at room temperature in the dark.
4.3 Results and Discussion
Figure 4.1 shows experimental and simulated RBS spectra for as-deposited ZTO
films. Peaks related to Zn and Sn are found at higher channel numbers in the spectra,
which corresponds to high backscattering energy, due to their high atomic masses.
Likewise, Si and O are found at lower channel numbers in the spectra, which correspond
to
ob
v
sp
1
(s
Ffr
o lower back
bserved due
alues of film
pectra [23]. T
.8. These va
The in
similar data w
igure 4.1: Rrom the ZTO
kscattering en
to the SiO2
m compositio
The RBS an
alues differ s
nset of Figur
were obtaine
RBS data fromO films after
nergy, due to
layer and th
on and areal
nalysis indica
lightly from
re 4.1 is an X
ed for as dep
m sputter-deannealing to
o their low a
he Si substrat
densities we
ates that the
m the specifie
XRD pattern
posited films
eposited ZTOo 600 °C.
atomic mass
te [22]. For t
ere varied to
as-deposited
ed nominal ta
n from ZTO f
s).
O films. The
. Two signal
the RBS sim
match the e
d film has a
arget compo
films anneal
inset shows
ls for Si are
mulations the
experimental
Zn/Sn ratio
osition.
led at 600 °C
s the XRD da
90
e
l
of
C
ata
91
The lack of strong diffraction peaks in the spectra indicates that the ZTO films are
amorphous even after annealing to these temperatures, which is consistent with prior
studies [24].
Figure 4.2 shows a depth profile from the Al/ZTO/Pt stack using ToF-SIMS. The
profile is plotted as the log of the ion intensity versus sputter depth time in seconds. The
left side of the figure corresponds to the top of the device whereas the right side of the
figure corresponds to bottom of the device. An AlOx interface layer was identified due to
the increase of the oxygen signal prior to an increase in intensity of the zinc or tin signals
from the ZTO film. Based on sputter etch rates we estimate that there was a 3-4 nm of
aluminum oxide at the interface between the Al top electrode and the ZTO film. It was
also found that the Al signal still has significant intensity ~1/3 of the way through the
ZTO layer, which can be contrasted by the relatively sharp interface at the Pt electrode
(although significant Zn intensity is observed in the Pt). This suggests that there may be
important interfacial reactions between Al and ZTO that take place due to the high
oxygen affinity of Al [25].
F
el
o
fo
o
in
sh
w
w
igure 4.2: To
Recen
lectrodes and
f the devices
ormation of
f metal botto
nfluence dev
Figure
hows an ima
was ground a
was applied t
oF-SIMS de
nt studies hav
d the oxide s
s [26]. Depe
a reduced la
om electrode
vice switchin
e 4.3 illustra
age of ZTO c
and voltage b
o the top ele
epth profile o
ve suggested
switching m
nding on the
ayer of ZTO
e and adhesio
ng [27].
ates the I-V c
cross bar stru
bias was app
ectrode, a lin
on Al/ZTO/P
d that the int
material can s
e metal top e
at the top in
on layer into
characteristic
uctures. For
plied on the t
near increase
Pt structure.
terfacial reac
strongly influ
electrode, thi
nterface regio
o the oxide f
cs of an Al/Z
all measure
top electrode
e in current w
ctions betwe
uence the I-V
is process ca
on. Likewise
film can sign
ZTO/Pt devi
ments the bo
e. When a ne
was obtained
een metal
V characteris
an result in t
e, the diffusi
nificantly
ice and the in
ottom electro
egative volta
d and at a cer
92
stics
the
ion
nset
ode
age
rtain
93
voltage there was a drastic increase in current as the device transitions from the initial
HRS to the LRS. This is the SET process (device switched “on”) and the voltage at which
this transition takes place is the SET voltage (VSET). In contrast, when a positive voltage
was applied to the top electrode, when the device is in the LRS, there was a transition
back to the HRS. This transition is the RESET process (device switched “off”), and the
voltage at which this transition takes place is the RESET voltage (VRESET).
To avoid electrical breakdown of the devices, the compliant current (CC) was
used during the SET process prior to performing the I-V measurements. We found that a
gentle forming process was required to obtain stable bipolar switching characteristics.
During the forming process, the limiting CC was slowly increased in steps from 100 nA
to 350 µA. We continued ramping up the CC until the devices switched from its initial
unipolar to bipolar switching. For these studies the forming voltage was relatively low
and the value was maintained close to the SET voltages for successive cycles to obtain
stable switching of the devices. Prior studies have also found that stable bipolar switching
could be obtained using low forming voltages and CCs of less than 1 mA [28] [29].
Figure 4.4 shows bipolar resistive switching for more than 30 cycles for a 20 x 20
µm2 Al/ZTO/Pt cross bar cell obtained after an initial gentle forming process. The CC
was set to 350 µA for these measurements. It can be seen that there was significant
variation in VSET while VRESET stayed fairly constant.
Fim
F
d
ar
w
el
el
ot
igure 4.3: Bmage of devi
The re
igure 4.4. W
ecreases wit
rea filament
with cell area
lectrode - ZT
lectrode area
ther materia
ipolar resistiice structure
esistance cha
We found that
th increasing
conduction
a the limiting
TO interface
a [8]. Simila
ls systems a
ive switchine.
ange (RLRS a
t RLRS was c
g cell area [3
is the propo
g conduction
e, or to non-c
ar scaling beh
as well [11][1
ng in ZTO RR
and RHRS) wi
constant with
0]. Typicall
osed conduct
n mechanism
contacting m
havior for R
15][17].
RAM crossb
ith the cell a
h increasing
ly when RLR
tion mechani
m is related to
multi-filamen
RLRS and RHR
bar device; in
area is plotte
cell area wh
RS doesn’t ch
ism. Since R
o interfacial
nt arrays that
RS has been o
nset shows a
d in the inse
hile RHRS
hange with c
RHRS does ch
effects at th
t scale with
observed for
94
an
et of
cell
hange
he
Fsc
I
b
su
re
d
P
I-
igure 4.4: Bchematic of
In ord
versus log V
oth LRS and
uggesting Oh
egion of HR
etermine the
oole-Frenke
-V data and t
ipolar resistithe device st
der to better u
V for the LR
d HRS (V <
hmic conduc
S (V > VRES
e current con
el and Schott
the followin
ive switchintack and dep
understand c
S and HRS a
VRESET) a lin
ction. As sho
SET) has two
nduction mec
tky emission
ng equation [
ng in ZTO RRpendence of
carrier transp
as shown in
near increase
own in the in
linear regim
chanism in t
n. Poole-Fren
[31],
RAM crossbRLRS and RH
port in these
Figure 4.5. A
e with a slop
nset of Figur
mes in the ln
the high volt
nkel emissio
bar device; in
HRS on cell a
e devices we
At the low v
pe of ~1 is ob
re 4.5, the hi
(I/V) versus
tage region w
on can be eva
nset shows aarea.
have plotted
voltage regio
bserved
igh voltage
s V1/2 plot. T
we consider
aluated using
95
a
d log
on of
To
both
g the
96
ln( / )~[( ) ] where q is electric charge, d is thickness of film, εr is dynamic dielectric constant and ε0 is
permittivity of free space, kB is Boltzmann’s constant and T is the absolute temperature.
From the slopes of ln (I/V) Vs V1/2, the values for εr can be estimated [31].
Also shown in the inset of Figure 4.5, the high voltage region of HRS (V >
VRESET) has two linear regimes in the ln (I) versus V1/2 plot. Schottky emission can be
evaluated with the I-V data and the following equation,
ln( )~[( ) ] 2
From the slopes of ln I Vs V1/2, the values for εr can also be estimated [32][33].
Table 4.1 gives the experimentally determined values for the optical dielectric constant
(εr) of ZTO, which was calculated using slopes (S1 and S2) assuming either Schottky or
Poole-Frenkel emission at the high voltage regions (V > VRESET) of HRS. For
comparison, the value of optical dielectric constant (εr,o) for ZTO was found to be ~ 4
using the refractive index (n=2) of the ZTO films obtained from ellipsometry and using
the relation n= εr,o1/2 [34][35].
Fthv
Tuhw
igure 4.5: Ohe insets shooltage region
Region o
S
S
Table 4.1: Exsing slopes (igh voltage r
was determin
Ohmic conduow plots for Sns (V > VRE
of analysis
S1
S2
xperimental v(S1 and S2) regions (V >
ned to be ~ 4
ction at low Schottky em
SET) of HRS
values for opof Schottky
> VRESET) of from spectr
voltage regimission and P
.
εr Schottk
0.78 ± 0.4
0.12 ± 0.
ptical dielectand Poole-FHRS. The o
roscopic ellip
ions of LRS Poole – Fren
ky
4
1
tric constantFrenkel condoptical dielecpsometry.
and HRS (Vnkel emission
εr Poo
5.6
0.5
t (εr) for ZTOduction mechctric constan
V < V RESET)n for the high
ole-Frenke
65 ± 2.5
55 ± 0.4
O calculatedhanism at th
nt εr,o for ZTO
97
) and h
l
d he O
98
As mentioned above the slope changes from one regime (S1) to another (S2) whether the
data was plotted ln (I) versus V1/2 or ln (I/V) versus V1/2, although S1 occurs over a wider
range of data. Changes in slope in the high voltage regime have also been observed in
prior studies [36]. For Schottky emission the calculated values (with T = 298 K and d =
50 nm) for εr were significantly smaller than εr,o for ZTO obtained from ellipsometry. For
Poole-Frenkel emission the value for εr obtained from region S1 gave a value fairly close
to εr,o for ZTO, however the value for εr obtained from region S2 was significantly lower
than εr,o. Since the S1 covers a wider range of data, we propose that Poole-Frenkel
emission was the dominant conduction mechanism for the high voltage regime of the
HRS. It has been shown in the literature that Poole-Frenkel emission dominates for oxide
films greater than 10 nm thick , whereas Schottky emission dominates for oxide films
less than 10 nm thick [33][37].
Figure 4.6 shows the statistical distribution of VSET and VRESET for more than 30
switching cycles. The values for mean of VSET and VRESET were determined to be -1.97
and 0.47 with standard deviations of 0.57 and 0.34 , respectively. These results suggest
that stable and uniform distribution for VSET and VRESET could be obtained for ZTO based
RRAM.
F
d
sw
gr
v
v
sh
sh
d
igure 4.6: St
enotes mean
Figure
witching cyc
reater than 1
olatility of th
oltage of 0.1
hown in Figu
The v
how any deg
ata retention
tatistical dist
n and σ deno
e 4.7 shows
cles. Despite
104 was obta
he ZTO RRA
1 V after var
ure 4.8.
alues for RH
gradation dur
n.
tribution of S
otes standard
the change i
e the fluctuat
ained which i
AM, retentio
rious lengths
HRS/RLRS wer
ring the tests
SET and RE
d deviation.
in RHRS and
tions of RHRS
is suitable fo
on measurem
s of time to o
e greater tha
s. This sugge
ESET voltage
RLRS with re
S, a resistanc
or most appl
ments were p
obtain the va
an 104 for lon
ests that the
es for 30 cyc
espect to the
ce ratio (RHR
lications. To
performed w
alues for RHR
nger than 10
ZTO RRAM
cles where µ
number of
RS/RLRS) of
test the non
with a read
RS and RLRS a
04 sec and did
M has excell
99
µ
n-
as
d not
lent
F
F0
igure 4.7: R
igure 4.8: R.1 V.
RLRS and RHR
Retention cha
S over 30 res
aracteristics o
sistive switc
of ZTO mem
ching cycles.
mory device
.
taken for a r
read voltage
100
e of
101
The proposed bipolar resistive switching model for the Al/ZTO/Pt cross bar
memory cell will be discussed below. The devices were initially in the HRS due to the
inclusion of oxygen during ZTO sputter deposition process, which limits the number of
oxygen vacancies, as well as reducing the carrier concentration in the ZTO. ToF-SIMS
suggests that an insulating AlOx interface layer was formed by a redox reaction between
Al and ZTO at the interface which is expected due to the enthalpies of formation for
Al2O3, ZnO, and SnO2 [38]. This interfacial redox reaction results in the formation of
oxygen vacancies and metal cation interstitials at the ZTO interface leading to a higher
carrier concentration.
During the forming process, when a negative voltage was applied to the Al top
electrode localized conductive filaments may be formed between the top and bottom
electrodes by the drift of oxygen vacancies (Vo2+) or metal interstitials (e.g., Zn2+, Al3+,
Sn4+, etc.) towards the top electrode, or by the drift of O2- ions towards the bottom
electrode. It may also be possible that the AlOx layer becomes more conductive due to the
migration of O2- ions from the AlOx layer and into the ZTO layer. It is the accumulation
of these defect species that forms the conductive filaments resulting in switching to the
LRS [14]. We also found that the device did not have stable resistive switching when the
top Al electrode was grounded and the voltage was applied to bottom electrode. This
suggests that the redox reaction at the Al/ZTO interface may be important for the
switching observed in our devices.
102
Alternatively the devices can be returned to the HRS by application of a positive
voltage at the Al electrode (RESET process). The switching to the HRS may be due to
reduction/oxidation reactions that cause the dissolution of filaments near the Al/ZTO
interface layer and the resulting drift of the defect species discussed above. It is also
possible that during the SET process (V<0), that devices switch from HRS to LRS due to
the transport of O2- ions from the AlOX interface layer into the ZTO films and likewise
for the RESET process (V>0) devices switch back to HRS due to the transport of O2- ions
from the ZTO films into the AlOX interface layer [39][40]. Further studies on the role of
Al intrinsic cation and oxygen vacancy drift need to be performed to better understand
the switching mechanisms.
4.4 Conclusion
In conclusion RRAM devices based on sputter deposited ZTO have been investigated for
non-volatile memory applications. The ZTO films were found to remain amorphous even
after annealing up to 600 °C with the Zn/Sn ratios of 1.8. The as-deposited amorphous
ZTO based RRAM devices had stable bipolar switching characteristics with a large
RHRS/RLRS ratio > 104 and the HRS and LRS were found to be stable for retention times >
104 sec. The driving mechanism behind resistive switching was proposed to be due to
combination of bulk effect (formation and dissolution of filamentary conduction paths)
and interface effect (redox reactions at the Al/ZTO interface) although further
investigations are required to determine the exact mechanism.
103
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107
CHAPTER 5 - CONCLUSIONS AND RECOMMENDATIONS FOR FUTURE WORK
5.1 Conclusions
In this thesis, TFTs with amorphous ZTO as a channel layer was demonstrated
successfully. The effect of both deposition and post annealing conditions on film
structure, composition, surface contamination, and TFT device performance were
studied. The XRD results proved that films were amorphous even after annealing at 600
°C. The RBS and XPS studies on the ZTO thin films indicated that the film stoichiometry
was close to the target stoichiometry of Zn2SnO4.The electrical test results on ZTO films
using TFT test structures indicated that mobilities as high as 17 cm2 V-1 s-1 and a high
value for drain current ON-OFF ratio could be obtained for depletion mode devices
suitable for practical applications. We also found that with changes in deposition
conditions and annealing temperature we can control the electrical properties of the films
where insulating, semiconducting, or conducting ZTO films could be obtained.
RRAM devices based on amorphous ZTO have been investigated for non-volatile
memory application. The as-deposited amorphous ZTO based RRAM devices exhibited
stable bipolar switching characteristics with a large RHRS/RLRS ratio > 104 which was
stable for long retention times of more than 104 sec The driving mechanism behind
resistive switching is proposed to be due to combination of bulk effect (formation and
rupture of filamentary conduction paths) and interface effect (redox reactions at the
108
Al/ZTO interface) although further investigations are required to determine the exact
mechanism.
5.2 Recommendations for future work
Though, we were able to achieve stable bipolar resistive switching with sputter
deposited amorphous ZTO material system, the mechanism behind the resistive switching
needs further investigation. A thorough understanding of the role of interface AlOx in
device switching is necessary to determine if the switching is due to changes at the
interface. For these studies, SIMS depth profiling can be performed with the devices at
various switching conditions and focus on how the signal of ions at the Al/ZTO and
ZTO/Pt interfaces are modified with different device conditions. In addition, conductive
atomic force microscopy (CAFM) or Kelvin probe force microscopy (KPFM) studies on
ZTO RRAM after removal of the top electrode can be used to determine if the driving
mechanism behind resistive switching supports conductive filament model or the
interface model for switching.
109
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124
Appendix: Sol-gel chemistry for zinc tin oxide (ZTO)
In order to make ZTO films by sol-gel method, zinc chloride ZnCl2 and tin oxide (SnO2)
powders were purchased from Alfa Aesar and used as starting materials. Since the SnO2
is very resistant to several solvents, SnO2 was converted to tin iodide which is also called
stannic iodide (SnI4) which readily dissolved in non-polar solvents. SnI4 is comparatively
expensive when purchased directly. So, concentrated hydriodic acid (HI) ACS grade,
(47%, stab. with 1.5% hypophosphorous acid) purchased from Alfa Aesar was used to
convert the SnO2 to SnI4. 0.1553 g of stannic oxide was added to the 3 ml of HI in a
reaction vessel with a magnetic stirrer. The reaction vessel was kept on a hot plate and at
about 90-100 °C, there was constant boiling of hydriodic acid and the reaction started.
The reaction vessel was closed with a watch glass. With the successive addition of HI
with specific time intervals the reaction was complete in less than 20 min. The end of
reaction was marked by formation of orange-red stannic iodide and on the walls of the
reaction vessel it appeared as a yellow to orange colored sublimate.
The reaction could be,
The excess HI along with non-dissolved tin oxide was separated from newly formed
stannic iodide by centrifugation. To remove water and excess acid the extracted SnI4
powder was dried in a desiccator filled with calcium chloride, metallic copper and
potassium hydroxide.
T
d
F
T
fo
pr
ad
th
P
5
The formatio
iffraction pa
igure A. 1: D
To make the
ormed SnI4
recursors ha
dded to the
he solution s
rior to spin-
min, washe
on of SnI4,
atters from S
Diffraction p
solution to
was dissolv
ad a mass ra
solution wh
uitable for s
-coating, the
d with aceto
was confirm
nI4.
patterns of ne
o prepare ZT
ved in 5 ml
atio of 2:1 f
hich acted as
pin coating.
substrates w
one, IPA and
med from X
ewly formed
TO films, 0
acetonitrile
for ZnCl2:Sn
s a stabilizer
The substra
were ultrason
DI water, bl
XRD measu
d SnI4.
0.05 M of Z
which acted
nI4. Few dr
r that helped
ates were SiO
nicated in a
low dried in
urements. F
ZnCl2 and 0
d as a comm
rops of ethy
d to increase
O2/Si.
solution of
n nitrogen an
Figure A1 s
0.05 M of n
mon solvent
ylene glycol
e the viscosi
0.1 M NaOH
nd heated on
125
shows
newly
. The
were
ity of
H for
a hot
p
sp
an
T
st
an
un
F
w
p
T
µ
F
T
an
late at 85 °C
pin rate of
nnealed at 6
The thicknes
tudies with t
nnealing up
nwanted ZT
inally the TF
with a devic
atterned by
The width/len
µm/100 µm.
igure A. 2: B
The TFTs we
nalyzer. For
C for 5 min. T
3500 rpm a
600 °C in a t
s of films w
the sputter d
to 600 °C.
TO was rem
FT devices w
ce stack as
depositing ~
ngth (W/L) r
Bottom gate
ere tested us
r these mea
The solution
and spin tim
tube furnace
was found to
eposited ZT
The ZTO w
moved using
were formed
shown in F
~500 nm of
ratios of the
TFT test str
sing probe st
asurements
n was finally
me of 30 se
e with a ram
o be 50 nm
O proved th
was patterne
oxalic acid
d using the so
Figure A2.
Al via therm
fabricated d
ructure.
tation with A
the drain c
y spin coated
ec to form t
mp rate of 10
confirmed f
hat ZTO film
ed using ne
(0.3 M) etc
olution proce
The source
mal evaporat
devices were
Agilent 415
current (ID)
d on SiO2/Si
thick films.
0 °C/min to
from ellipso
ms were amor
gative photo
ch which to
essed ZTO a
e and drain
tion through
e 1000 µm/2
5C semicon
and gate c
substrates w
The films
form ZTO f
ometry. The
rphous even
olithography
ook about 4
as a channel
electrodes
h a shadow m
200 µm and
nductor param
current (IG)
126
with a
were
films.
prior
n after
y and
min.
layer
were
mask.
1000
meter
were
m
d
T
ch
m
F
O
is
A
measured wh
ouble sweep
The ZTO fi
haracteristic
mobility (µavg
igure A. 3: T
One recomm
sotopically e
A4 shows the
hile the gate
p mode whe
films were
s of the TF
g) was found
Transfer cha
mended futur
enriched Zn
e device stac
to source v
ere the drain
semiconduc
FT test struc
d to be ~ 0.15
aracteristics o
re work cou
and Sn prec
ck that can be
voltage (VGS
n to source v
cting which
cture as show
5 cm2/V-s.
of the solutio
uld be using
cursors to pe
e used for di
S) was scann
voltage (VD
h can be p
wn in Figur
on processed
g solution pr
erform diffu
iffusion stud
ned from -20
S) was held
proved from
re A3. The
d ZTO TFT.
rocessed ZT
sion studies
dies.
0 V to 20 V
constant at
m the ON
value of av
TO formed u
in SIMS. F
127
V in a
1 V.
-OFF
erage
using
Figure
F
T
su
is
aw
is
co
io
igure A. 4: D
The as-depos
ubjected to a
sotopically l
way from th
sotopes, by p
oncentration
ons (O2-) or m
Device struc
sited films c
annealing un
abeled 68Zn,
he sputtered
performing a
n of Zn, Sn
metallic ions
cture for diffu
can be stack
nder labeled
, 112Sn or 18O
d ZTO film
a depth prof
and O speci
s (m+) cause
fusion studie
ked with the
oxygen (18O
O ions from
ms. In SIMS
file analysis
ies can help
s the resistiv
es using SIM
e above stru
O). This ann
m the solution
is very sen
on this dev
p to confirm
ve switching
MS.
ucture and t
nealing can c
n processed
nsitive towa
vice stack an
m if the mov
g in RRAM.
then they ca
cause diffusi
d ZTO towar
ards detectio
nd monitorin
ement of ox
128
an be
ion of
rds or
on of
ng the
xygen
