Electrical Memory Characteristics of Nitrogen-Linked Poly(2,7

ARTICLE pubs.acs.org/JPCC

Electrical Memory Characteristics of Nitrogen-Linked Poly(2,7-carbazole)s Suk Gyu Hahm,†,§ Taek Joon Lee,†,§ Dong Min Kim,†,§ Wonsang Kwon,† Yong-Gi Ko,† Tsuyoshi Michinobu,*,‡ and Moonhor Ree*,† †

Department of Chemistry, Division of Advanced Materials Science, Center for Electro-Photo Behaviors in Advanced Molecular Systems, Pohang Accelerator Laboratory, Polymer Research Institute, and BK School of Molecular Science, Pohang University of Science & Technology, Pohang 790-784, Republic of Korea ‡ Global Edge Institute, Tokyo Institute of Technology, 2-12-1 Ookayama, Meguro-ku, Tokyo 152-8550, Japan

bS Supporting Information ABSTRACT: We studied the electrical memory characteristics of the following nitrogen-linked poly(2,7-carbazole)s: poly(9hexadecyl-2,7-carbazole-alt-N,N-(4-hexadecyloxy)aniline), poly(9-hexadecyl-N,N0 -diphenylcarbazole-2,7-diamine-alt-1,3benzene), and poly(9-hexadecyl-N,N0 -diphenylcarbazole-2,7diamine-alt-4,40 -biphenyl). These polymers are amorphous; however, in thin films, they are slightly oriented in the film plane. All polymers in devices with aluminum top and bottom electrodes were found to exhibit similar dynamic random access memory (DRAM) behaviors without polarity. They are operable with a low voltage (less than (3 V) and a high ON/OFF current ratio (105
109, depending on the polymer) over the thickness range 8
60 nm. The memory behaviors were found to be governed by space-charge limited conduction and local filament formation. These memory characteristics might originate from the electrondonating carbazole and triphenylamino units in the polymer backbones, which act as charge-trapping sites but have weak electric polarization because of the absence of counterparts. Overall, these polymers are suitable active materials for the mass production at low cost of high-performance, programmable volatile memory devices.

’ INTRODUCTION High-performance memory devices are in high demand because of the rapid growth of the market for mobile devices.1 Currently, all memory devices are fabricated with inorganic materials. Inorganic materials require elaborate fabrication processes, such as vacuum evaporation and deposition. Moreover, these devices have some general shortcomings (limited endurance, high voltage consumption, etc.). In contrast, organic materials can easily be miniaturized, and their properties can easily be tailored through chemical synthesis. As a result, their use as alternative materials for memory devices has been considered.2 However, organic materials also have some disadvantages, such as elaborate fabrication processes (vacuum evaporation and deposition), low boiling points, and low chemical resistances. Polymer materials have significant advantages over inorganic materials as well as organic small molecules. They exhibit easy processability, flexibility, high mechanical strength, and good scalability. Moreover, they can be processed at low cost, and with their use the multistack layer structures required for highly dense memory devices can easily be fabricated. Therefore, there is significant current research into the development of polymer materials that meet the requirements of memory devices. As a result, some candidate polymer materials have been reported.3
7 r 2011 American Chemical Society

In particular, several polymer systems containing the carbazole moiety have been reported to exhibit nonvolatile memory characteristics as active layers in electrical memory devices.8,9 Furthermore, carbazole-containing polymers have been tested as hole-transport materials in organic light emitting diodes (OLEDs), organic field effect transistors (OFETs), and organic solar cells.10 Some polymers containing diphenylamino and triphenylamino moieties were found to exhibit volatile or nonvolatile memory characteristics as active layers in such devices.11
13 Recently, a series of nitrogen-linked poly(2,7-carbazole) polymers were introduced and characterized as OLED materials:14 poly(9hexadecyl-2,7-carbazole-alt-N,N-(4-hexadecyloxy)aniline) (P1), poly(9-hexadecyl-N,N0 -diphenylcarbazole-2,7-diaminealt-1,3-benzene) (P2), and poly(9-hexadecyl-N,N0 -diphenylcarbazole-2,7-diamine-alt-4,40 -biphenyl) (P3) (Figure 1a). The nitrogen-linked poly(2,7-carbazole) polymers (P1, P2, and P3) consist of triphenylamino moieties in the backbone in addition to the carbazole units. Interestingly, the triphenylamino moieties are chemically linked with the carbazole units in a hybrid manner (Figure 1a). Moreover, these polymers exhibit polyimide-like Received: July 28, 2011 Revised: October 1, 2011 Published: October 05, 2011 21954

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These polymers are easily fabricated by means of conventional solution spin-coating (or roll- or dip-coating) and subsequent drying, producing high-quality nanoscale thin films that are suitable as active polymer memory layers. In the fabrication of the memory devices, metal electrodes were prepared by electronbeam (E-beam) deposition or thermal evaporation. The devices were found to exhibit excellent dynamic random access memory (DRAM) characteristics with high ON/OFF current ratios. The DRAM characteristics of the devices were found to vary with the thickness of the active polymer layer. In addition, the interfaces of the active polymer layers with the substrates, metal electrodes, and air were examined in detail by using synchrotron X-ray reflectivity (XR). The structures of the polymers in thin films were further investigated with X-ray scattering. The mechanism of the DRAM behavior was also investigated.

Figure 1. (a) Chemical structures of nitrogen-linked poly(2,7-carbazole) polymers. (b) Schematic diagram of the memory devices fabricated with nanometer-scale thin films of the polymers and Al top and bottom electrodes. (c) Polymer thin films in contact with silicon substrate, Al bottom electrode, Au bottom electrode, and Al top electrode.

thermal and mechanical stability.14 Considering their chemical structures, properties, and OLED performance, the nitrogenlinked poly(2,7-carbazole) polymers may be promising candidates for active materials in electrical memory devices. In this study, we further tested the nitrogen-linked poly(2,7carbazole) polymers as active layer materials in memory devices.

’ EXPERIMENTAL SECTION The nitrogen-linked poly(2,7-carbazole)s (Figure 1a) were synthesized according to a previously reported synthetic method.14 For the fabrication of memory devices, homogeneous solutions of the nitrogen-linked poly(2,7-carbazole)s were prepared with a concentration of 0.5 to 2 wt % in cyclopentanone and filtered through PTFE-membrane microfilters with a pore size of 1.0 μm. Nanoscale thin films of 8, 15, 35, 60, and 110 nm thickness were prepared from the filtered solutions by spincoating onto silicon substrates and bottom electrodes at 2000 rpm for 60 s, followed by drying on a hot plate at 80 °C in vacuum for 8 h. The resulting films’ thickness was determined by using a spectroscopic ellipsometer (model M2000, Woollam). Optical properties were measured with a Scinco ultraviolet
visible (UV
vis) spectrometer (model S-3100). Cyclic voltammetry (CV) measurements were performed in a 0.1 M solution of tetrabutylammonium tetrafluoroborate in acetonitrile by using an electrochemical workstation (IM6ex impedance analyzer) with a platinum gauze counter electrode and an Ag/AgCl (3.8 M KCl) reference electrode, and polymers were coated onto the Au bottom electrode, which was deposited on a silicon wafer. A scan rate of 100 mV/s was used. Grazing-incidence X-ray scattering (GIXS) measurements were performed at the 4C1 and 4C2 beamlines15,16 of the Pohang Accelerator Laboratory (PAL), Pohang University of Science and Technology (POSTECH). Measurements were performed at a sample-to-detector distance (SDD) of 125 mm for grazing incidence wide-angle X-ray scattering (GIWAXS) and 2220 mm for grazing incidence small-angle X-ray scattering (GISAXS). Scattering data were typically collected for 30 s using an X-ray radiation source of λ = 0.138 nm with a 2D charge-coupled detector (CCD: Roper Scientific, Trenton, NJ). The incidence angle αi of the X-ray beam was set at 0.147° for wide-angle scattering and 0.159° for small-angle scattering, respectively, which is between the critical angles of the polymer film and the silicon substrate (αc,f and αc,s). Scattering angles were corrected according to the positions of the X-ray beams reflected from the silicon substrate with respect to a precalibrated silver behenate (TCI, Japan) powder. Aluminum foil pieces were applied as a semitransparent beam stop because the intensity of the specular reflection from the substrate was much stronger than the scattering intensity of the polymer films near the critical angle. For the fabrication of devices (Figure 1b), aluminum (Al) was deposited on glass substrates as a bottom electrode with a thickness of 300 nm by electron beam sputtering, whereas Al 21955

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Table 1. Molecular Weights and Thermal Properties of the Polymers polymer

Mwa

Mnb

Mw/Mnb

P1

10 200

6300

1.61

P2 P3

26 000 75 600

11 700 26 400

2.22 2.86

Tg (°C)c

Td (°C)d

138 116

440 425

360

a

Determined by careful integration of proton nuclear magnetic resonance spectra. b Determined by gel permeation chromatography (tetrahydrofuran eluent, calibrated by polystyrene standards). c Determined by the second heating scan of differential scanning calorimetry measurements in nitrogen ambient. d Temperature at which 5% weight loss occurred in heating under nitrogen ambient.

top electrodes were thermally evaporated at a pressure of ca. 10
7 Torr through a shadow mask. The top electrodes had a thickness of 300 nm, and their area ranged from 2.0  2.0 to 0.5  0.5 mm2. Current
voltage (I
V) measurements and stress tests of the device were carried out in air ambient using a Keithley 4200-SCS semiconductor characterization system. I
V curves were recorded by performing forward and reverse voltage scans between
4.0 and +4.0 V at a scan rate of 500 mV/s. For XR experiments, we additionally prepared three different types of samples for each polymer, as shown in Figure 1c. In these samples, the polymer layer was targeted to have a thickness of 35 nm, whereas the metal electrode layer was prepared with a thickness of 10 nm. The surface roughness and morphology were also examined using an atomic force microscope (Multimode AFM Nanoscope IIIa, Digital Instruments). XR data were measured at the 3C2 and 8C1 beamlines17 of PAL. Samples were mounted on a Huber four-circle goniometer equipped with a scintillation counter with an enhanced dynamic range (Bede Scientific, EDR) as a detector. The X-ray beam with a wavelength λ of 0.1541 nm was collimated at the sample position to 2 mm (horizontal) by 0.1 mm (vertical). Specular reflection was measured in θ
2θ scanning mode. The reflectivity, R, that is, the ratio of the reflected beam intensity to the primary beam intensity, was measured; here the primary beam intensity was monitored with an ionization chamber.

’ RESULTS AND DISCUSSION The nitrogen-linked poly(2,7-carbazole) polymers (P1, P2, and P3) were synthesized with a weight-average molecular weight Mw of 10 200
75 600 and a polydispersity index of Mw/Mn (Mn is the number-average molecular weight) of 1.61 to 2.86 (Table 1). The polymers are soluble in common organic solvents such as tetrahydrofuran, methylene chloride, and cyclopentanone. However, P1 was found to reveal relatively higher solubility than the other polymers; moreover, the solubility of P2 is relatively higher than that of P3. In general, aliphatic groups are more soluble in organic solvents than aromatic groups. Therefore, the observed solubility differences might be attributed to the differences in the contents of n-hexadecyl group per repeat unit. The content of such n-alkyl group per the unit length of the polymer backbone is in the increasing order P3 < P2 < P1 (Figure 1a). The glass-transition temperature, Tg, was found to be 138 °C for P2 and 116 °C for P3; P1 does not undergo a glass transition below its degradation temperature (Table 1). The degradation temperature Td in nitrogen ambient was determined to be 360 °C for P1, 440 °C for P2, and 425 °C for P3 (Table 1).

Figure 2. Representative synchrotron GIXS patterns of P1 in thin films (35 nm thick) coated on silicon substrates: (a) 2D GISAXS pattern measured at 25 °C with αi = 0.159° and (b) 2D GIWAXS pattern measured at 25 °C with αi = 0.147°.

Considering these polymers’ chemical structure (Figure 1a), the rigidity of backbone component is in the increasing order phenyl < biphenyl < carbazole. The fraction of the relatively most rigid carbazole unit per repeat unit is in the increasing order P3 < P2 < P1. Moreover, the polymers have two to three kink points per repeat unit. Because of the kink points and their interdistances, the overall cross-sectional diameter of the polymer chain is estimated to be in the increasing order P2 < P1 < P3. In general, the n-alkyl group is thermally less stable than aromatic groups and their derivatives. Taking these characteristics into account, the differences in the thermal properties of the polymers are attributed to the differences in their chemical structures. Regarding the relatively highest fraction of the rigid carbazole units in the backbone, P1 is expected to have higher Tg than those of P2 and P3. However, P1 exhibited the lowest Td due to the relatively highest content of n-hexadecyl groups per repeat unit. The highest occupied molecular orbital (HOMO) and lowest unoccupied molecular orbital (LUMO) energy levels of the P1, P2, and P3 polymers were determined by using UV
vis spectroscopy and CV analysis. (See Figures S1 and S2 in the Supporting Information.) All polymers in thin films showed the longest absorption maximum λmax in the range of 380
420 nm (Figure S1 in the Supporting Information). The λmax value is in the increasing order P2 < P3 < P1. These results indicate that P1 has the longest effective conjugation length due to the carbazole units in the backbone, P3 has the intermediate effective conjugation length due to the carbazole and biphenyl units, and P2 has the shortest effective conjugation length due to the carbazole and m-phenyl units. From the UV
vis spectra, the optical band gaps were determined to be 2.76 eV for P1, 3.04 eV for P2, and 2.90 eV for P3. The onset potentials for the first oxidation were observed at 0.76, 0.69, and 0.95 V for P1, P2, and P3, respectively (Figure S2 in the Supporting Information). The polymers showed somewhat different onset potentials for the first oxidation process. These results suggest that the differences in the chemical components of backbone (i.e., comonomer structures) affect the energy levels of the resulting polymers. On the basis of these oxidation process potentials, the HOMO levels of the polymers were estimated. In addition, the LUMO levels were 21956

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Figure 3. AFM surface images: (a) P1 film (35 nm thick) coated on silicon substrate; (b) Al electrode (10 nm thick) deposited on silicon substrate; (c) P1 film (35 nm thick) coated on the Al electrode deposited on silicon substrate; and (d) Al top electrode deposited on the P1 film supported with silicon substrate.

calculated from the HOMO levels and the optical band gaps of the polymers. The determined HOMO and LUMO levels were
5.02 and
2.26 eV for P1,
4.95 and
1.91 eV for P2, and
5.18 and
2.28 eV for P3, respectively. Morphological structures of the polymers in nanoscale thin films deposited on silicon substrates were examined with synchrotron GIXS analysis with a grazing incidence angle αi of 0.147
0.159°. It is noted that for the GIXS measurements with such αi values, X-ray beam can penetrate into the whole polymer film layer. The P1 polymer films exhibited a featureless GISAXS pattern (Figure 2a), which indicates that there is no discernible nanostructure or microstructure in the film. Furthermore, the P1 showed GIWAXS pattern that contains only a broad, weak scattering ring like an amorphous halo (Figure 2b). The scattering ring was determined to have a d spacing of 0.480 nm, which corresponds to the mean interchain distance of the polymer molecules. However, the scattering rings of the polymer films are very slightly anisotropic toward the meridian line. This scattering result indicates that the P1 polymer molecules in the film orient somewhat preferentially in the film plane rather than randomly. Similar GISAXS and GIWAXS patterns were observed for the P2 and P3 polymer films. (See Figure S3 in the Supporting Information.) The film surfaces of the polymers were examined by using atomic force microscopy (AFM). The P1 polymer films coated on silicon substrates were determined to have a root-meansquare (rms) roughness of 0.25 nm over an area of 1.0  1.0 μm2 (Figure 3a). In comparison, the P2 and P3 polymer films were found to have rms roughnesses of 0.93 and 1.72 nm, respectively (data not shown). In the case of the films coated on the Al bottom electrodes supported with Si wafers, the rms surface roughnesses were 0.26 nm for P1 (Figure 3c), 0.80 nm for P2, and 1.89 nm for P3. The Al bottom electrodes were found to have an rms surface roughness of 0.42 nm (Figure 3b), whereas the rms surface roughnesses of the Al electrodes deposited onto the polymer films were 0.68 nm for P1 (Figure 3d), 0.99 nm for P2, and 2.43 nm for P3. These results collectively indicate that the polymer of better solubility in organic solvents produces

Figure 4. XR profiles of the nitrogen-linked poly(2,7-carbazole) polymer films (ca. 35 nm thick) in contact with silicon substrate, Al bottom electrode, and Al top electrode (ca. 10 nm thick) (Figure 1c). The symbols are the measured data and the solid line represents the fit curve assuming a homogeneous electron density distribution within the film. The inset shows a magnification of the region around the two critical angles: αc,f and αc,s are the critical angles of the film and the substrate (silicon substrate or Al electrode), respectively.

nanoscale thin films of smoother surface via solution coating and subsequent drying process. Overall, all polymers gave reasonably good-quality films. In the devices, the polymers are in physical contact with metal electrodes. Therefore, it is valuable to examine their interfaces to interpret the devices’ performance. The polymer interfaces with the silicon substrates and metal electrodes were investigated with synchrotron XR analysis. The measured XR data are shown in Figure 4. All XR data could be satisfactorily fitted with Parratt’s fitting algorithm18 (Figure 4). From these XR analyses, the electron density profiles across the film thickness and other structural parameters were obtained. The results are summarized in Figure 5 and Table 2. Figure 5 shows the whole electron density profiles of the Al/polymer/Al devices, which were produced by combining the XR data analysis results of the Al/ polymer and polymer/Al samples. For the polymer films deposited onto the Al bottom electrodes, the XR analysis determined that the electron density Fe is 337.0 nm
3 for the P1 film, 363.4 nm
3 for the P2 film, and 397.1 nm
3 for the P3 film (Table 2). For each polymer, the determined Fe value is reasonably well-consistent with those measured for the polymer films deposited on silicon substrates and Au electrodes (Table 2). The XR analysis also found that an 21957

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aluminum oxide layer was formed with a thickness of 0.6 to 1.2 nm. Such aluminum oxide layer in contact with the polymer

Figure 5. Electron density profiles of the nitrogen-linked poly(2,7carbazole) polymer films (ca. 35 nm thick) in contact with silicon substrate, Al bottom electrode, and Al top electrode, which were obtained by analysis of the XR data (Figure 4).

film was determined to have an rms interfacial roughness of 1.1 to 1.2 nm, depending on the polymer (Figure 5 and Table 2). These interfacial roughness values are comparable to those determined for the polymer films coated on silicon substrates and Au electrodes (Table 2). These results collectively indicate that no aluminum atoms or ions were diffused into the polymer film layers during the process of polymer film deposition onto the Al bottom electrode. In addition, the XR analysis found that for the Al top electrodes deposited onto the polymer films, an aluminum oxide layer is present with a thickness of 0.5 to 1.0 nm, depending on the polymer (Figure 5 and Table 2). The rms interfacial roughness of the aluminum oxide layer in contact with the polymer film was determined to be 1.2 to 1.7 nm, depending the polymer (Table 2). The obtained interfacial roughness values are comparable to those determined for the polymer films coated on silicon substrates and Al and Au electrodes (Table 2). Moreover, the polymer films’ electron densities were retained without any changes due to the Al top electrode deposition process. These results indicate that no aluminum atoms or ions were diffused into the polymer film layers during the deposition process of Al top electrode onto the polymer film. Here the above determined electron densities Fe of the polymer films are discussed further with taking into account their chemical structure. The Fe values are in the increasing order P1 < P2 < P3. These results indicate that even though all polymers are amorphous, their chemical structure is directly reflected into the film density. In particular, the n-hexadecyl side group may disturb the lateral packing of the polymer chains. Therefore, higher content of n-hexadecyl groups per backbone unit length leads to lower electron density in the polymer film. We fabricated Al/polymer/Al devices and investigated their electrical memory characteristics. Figure 6 shows typical I
V characteristics of the 35 nm thick P1, P2, and P3 films with Al top and bottom electrodes, which were measured in dual sweep mode with a positive and negative bias (0 V f (4.0 V f 0 V) by using a semiconductor parameter analyzer with a compliance current of 0.1 A. Initially, the P1 film in the device is in the OFF state, which has a current level that is quite low (on the order of 10
14 to 10
10 A).

Table 2. Structural Parameters and Electron Density Profiles of Various Bilayer Samples Prepared from the Polymer Films, Silicon Substrates, Al Electrodes and Au Electrodes substrate

polymer

interlayera

top layer

sample (top/bottom) d (nm)b Fe (nm
3)c σ (nm)d d (nm)b Fe (nm
3)c σ (nm)d d (nm)b Fe (nm
3)c σ (nm)d d (nm)b Fe (nm
3)c σ (nm)d 692.5

0.1

38.4

338.9

1.0

6.9

676.2

0.7

P1/Al

12.9

865.9

1.0

34.7

337.0

0.8

1.2

1123.4

1.1

P1/Au

10.1

3985.8

0.9

33.0

337.3

0.5

34.0

339.0

1.0

P1/Si

Al/P1 P2/Si P2/Al P2/Au

9.5

692.9 861.9

1.4 1.6

38.1 37.7.

362.0. 363.4

0.8 0.7

10.1

4180.8

0.7

36.0

360.9

0.9

32.1

367.4

1.1

Al/P2 P3/Si

528.7

1.2

678.3 1095.7

1.0 1.1 1.3

0.5

548.6

397.5

1.8

10.92

677.3

0.4

0.6

1135.1

1.2

1.0

650.3

1.7

1.1

39.9

397.1

2.4

P3/Au

10.5

4024.8

0.8

35.2

398.4

1.3

39.3

399.3

1.3

9.7

802.8

1.2

0.7 11.2 1.0

39.5

794.3

807.4

1.2

0.5

9.5

9.4

796.4

696.1

P3/Al Al/P3

10.4

1.8

a

Silicon oxide layer for Si(bottom)/polymer(top), aluminum oxide layer for Al/polymer, and polymer-Al mixed layer (which is due to the roughness of interface) for polymer/Al systems. b Layer thickness. c Electron density of layer. d Roughness of layer in contact with air, lower or upper layer. 21958

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Figure 7. Experimental and fitted I
V curves for the Al/polymer (35 nm thick)/Al device: (a) OFF state with a combination of the Ohmic current model and the space-charge-limited conduction model and (b) ON state with the Ohmic model.

Figure 6. Typical I
V curves of the Al/polymer (35 nm thick)/Al devices, which were measured with a compliance current set of 0.1 A in dual sweep mode. The applied voltage was swept from 0 V f ( 4.0 V f 0 V. The electrode contact area was 0.5  0.5 mm2.

In the first dual sweep (see the first sweep in Figure 6a), there is an abrupt increase in the current (from 10
10 to 10
1 A) at 2.22 V (the switching-ON threshold voltage), which indicates that the film undergoes an electrical transition from an OFF state to an ON state. This electrical transition can serve as the “writing” process in a memory device. As soon as the first sweep is completed, the current returns to the OFF state; this transition can serve as the “erasing” process in the memory device. The erased state can again be written to a stored state when a voltage greater than the switching-ON threshold voltage is applied (see the second sweep in Figure 6a), which indicates that this memory device is rewritable. The third sweep was carried out after turning off the power for 1
3 s. Then, it was found that the

ON-state had already relaxed to the steady OFF-state. However, the film can be further programmed to the ON-state. The short retention time of the ON-state indicates that the memory device is volatile. The above processes can be repeated many times for every cell (Figure 6a). Overall, the 35 nm thick P1 films exhibit interesting and unique DRAM characteristics during positive voltage sweeps. For the negative bias sweeps (0 V f
4.0 V f 0 V), the DRAM memory characteristics were found to be the same as those for the positive bias sweeps (Figure 6a). In the Figure, the switching-ON threshold voltage under negative bias ranges from
2.25 to
2.65 V for the fourth, fifth, and sixth sweeps. The P2 and P3 polymers exhibit switching behaviors that are similar to those of the P1 polymer (Figure 6). However, the OFF-state current level and the switching-ON voltage are somewhat dependent on the polymer. To explain these memory behaviors, we further analyzed the measured I
V characteristics were further in detail with various conduction models. The Ohmic contact model was found to fit satisfactorily the I
V data for the ON state (Figure 7a). These results indicate that Ohmic conduction is dominant for all polymers in the ON state. Moreover, the current levels of the devices in the ON state were found to be independent of the device cell size, which is indicative of heterogeneously local filament formation. 6,19 The trap-limited space-charge limited conduction (SCLC) model was found to fit satisfactorily the I
V data for the OFF state (Figure 7b). The logarithmic plots of the I
V data for the OFF-state contain two linear regions for