Thermally-Driven Resistive Switching in Solution Processable Thin

Letter Cite This: J. Phys. Chem. Lett. 2017, 8, 5008-5014

pubs.acs.org/JPCL

Thermally Driven Resistive Switching in Solution-Processable Thin Films of Coordination Polymers Shammi Rana,† Anupam Prasoon,† Plawan Kumar Jha,† Anil Prathamshetti,‡ and Nirmalya Ballav*,† †

Department of Chemistry, Indian Institute of Science Education and Research (IISER), Dr. Homi Bhabha Road, Pashan, Pune 411 008, India ‡ Department of Physics, Indian Institute of Science Education and Research (IISER), Dr. Homi Bhabha Road, Pashan, Pune 411 008, India S Supporting Information *

ABSTRACT: Metal−organic coordination polymers (CPs) downsized to thin films with controllable electrical conductivity are promising for electronic device applications. Here we demonstrate, for the first time, thermally driven resistive switching in thin films of semiconducting CPs consisting of silver ion and tetracyanoquinodimethane ligand (AgTCNQ). High-quality and highly hydrophobic thin films of Ag-TCNQ were fabricated through a layer-by-layer approach upon sacrificing a predeposited layer of Cu-TCNQ on a thiolated Au substrate. Reversible switching between the high-resistance state (HRS) at 300 K and the low-resistance state (LRS) at 400 K with an enhancement factor of as high as ∼106 in the electrical resistance was realized. The phenomenon is attributed to the alternation of the Schottky barrier at the metal−semiconductor interface by thermal energy and not due to the formation of a conductive filament. Our discovery of thermally driven resistive switching as well as sacrificial growth of CP thin films on an organically modified substrate holds promise for the development of solution-processable nonvolatile memory devices.

M

(TCNQ) significantly enhanced the electrical conductivity of thin films of a well-known CP comprised of a Cu(II) ion and benzene tricarboxylic acid (BTC) ligand (HKUST-1).5,11,14 Recently, successful growth of semiconducting Cu-TCNQ thin films on various substrates other than Cu has also been achieved by LbL, and an unusual electrical rectification was observed upon exposing Cu-TCNQ thin films to iodine (I2) vapor.15 Liquid-phase heteroepitaxy (LPH) was adopted to empower LbL in sequentially growing distinctive CPs on a single-layered structure.6,16,17 LPH enabled fabrication of several bilayered (also trilayered) thin films of CPs with similar as well as different lattice constants.6,18,19 Of particular mention are the heterostructured thin films of Cu-based CPs with distinctive space groups (tetragonal and face-centered cubic).18,19 Herein, we present an example of sacrificial growth of thin films involving CPs (Figure 1). Despite having similar lattice parameters, LPH did not produce a bilayered/hybrid thin film and the seeding layer acted as a sacrificial layer to grow the other layer on top of the SAM in controlled manner. Specifically, we have grown a homogeneous Ag-TCNQ thin film upon a sacrificing Cu-TCNQ thin film. Notably, the LbL approach did not produce thin films of Ag-TCNQ directly likewise, our recently developed Cu-TCNQ thin films.15 Also,

etal−organic materials with controllable electronic and magnetic properties are useful for thin film device applications. On the one hand, flexible design of metal and organic constituents allows fine-tuning of the electronic and magnetic properties in these materials. On the other hand, electronic and magnetic properties of such materials could be susceptible to physical and chemical perturbations, for example, by light, electric field, heat, and chemical stimuli.1,2 Thus, metal−organic materials with bistable electronic and magnetic characteristics are desirable to configure the devices easily. Metal−organic coordination polymers (CPs) have emerged as an important class of functional solid-state metal−organic materials in the past 2 decades.3,4 To explore various CPs in the domain of technological applications, development of highquality thin films is essential.5−7 An elegant approach of growing thin films of CPs on organically functionalized surfaces is liquid-phase epitaxy (LPE) or a layer-by-layer (LbL) approach where self-assembled monolayers (SAMs), for example, thiolates on a Au substrate, can be used as the seeding templates.8,9 LbL offers unique opportunities: (i) highly oriented and uniform crystal growth, (ii) easily controllable thickness, (iii) minimizing the interpenetration of CPs,10 and (iv) finally, operational simplicity (ambient conditions and a step-by-step method). However, the LbL approach is found to be limited to either a Cu(II) or a Zn(II) ion, and also, the thin films of CPs are observed to be poor electrical conductors.5,8,11−13 Infiltration with specific redoxactive species such as ferrocene and tetracyanoquinodimethane © 2017 American Chemical Society

Received: August 14, 2017 Accepted: September 25, 2017 Published: September 25, 2017 5008

DOI: 10.1021/acs.jpclett.7b02138 J. Phys. Chem. Lett. 2017, 8, 5008−5014

Letter

The Journal of Physical Chemistry Letters

In step I, we grew Cu-TCNQ thin films through the LbL approach, and subsequently, Cu-TCNQ thin films were first dipped into AgNO3 solution and then dipped into TCNQ solution to make 1 cycle of the LbL growth, and 10 consecutive cycles were carried out (step II). A remarkable change in the color of the resulting thin film was observed; the initial dark green color of Cu-TCNQ was completely turned into dark blue color and suggested almost exclusive formation of Ag-TCNQ (Figure 2a). Despite having different chemical identity, the highly hydrophobic characteristic of the initial surface15 was also retained; contact angle (CA) values of H2O on Ag-TCNQ thin films were measured to be ∼130° (Figure 2a). Growth of the thin film was duly monitored in each cycle by recording powder out-of-plane X-ray diffraction (PXRD) patterns starting from 0th to 10th cycle (Figure 2b). In the PXRD pattern of the initial thin film (0th cycle), peaks at 2θ values of 11.0, 15.6, and 17.4° are characteristics of Cu-TCNQ material corresponding to (011), (002), and (012) planes of the phase-I structure, respectively.20,21 With progressive cycles, the intensities of these peaks were gradually decreased from the 0th cycle and finally vanished in the 10th cycle. Instead, new diffraction peaks at 2θ values of 10.25, 14.5, and 21.5°, which are characteristic of Ag-TCNQ, appeared, corresponding to (002), (022), and (004) planes of the phase-I structure, respectively.21,22 In energy dispersive X-ray spectroscopy (EDXS) data, the characteristic Cu Lα line at ∼0.93 keV was exclusively visible in the initial film; however, in the final thin film, the Cu Lα line was found to be almost absent, and characteristic Ag Lα (∼2.93 keV) and Ag Lβ (∼3.15 keV) lines appeared (Figure 2c).23 Elemental mapping revealed no trace of Cu and almost an exclusive presence of Ag, uniformly across the final thin film

Figure 1. Schematic diagram of sacrificial LbL growth. Dipping of a preformed Cu-TCNQ thin film into AgNO3 solution (1), into TCNQ solution (2), and again into AgNO3 solution (3). Successive dipping into AgNO3 and TCNQ solutions makes one cycle of LbL. Cu-TCNQ grown on a SAM/Au surface by LbL is expected to act as a seeding layer for the subsequent growth of Ag-TCNQ (4). Surprisingly, a thin film of Ag-TCNQ (4a) only was observed instead of hybrid AgTCNQ/Cu-TCNQ (4b), whereby the initial Cu-TCNQ thin film was sacrificed.

changing the terminal groups of SAMs from −COOH to −NH2 or −OH/−SH (Figure S1) as well as the growth temperature did not help to grow Ag-TCNQ crystals.

Figure 2. (a) Optical images of Cu-TCNQ (green) and Ag-TCNQ (blue) thin films. Optical image of a distilled water droplet on the Ag-TCNQ thin film surface. (b) Out-of-plane XRD patterns showing gradual transformation of Cu-TCNQ to Ag-TCNQ in thin film configurations. (c) EDXS spectra recorded on thin films obtained in step I (Cu-TCNQ) and step II (Ag-TCNQ). (d) Elemental mapping of C, N, Cu, and Ag on the thin film obtained in step II. (e) Ag 3d and Cu 2p XPS data recorded on the thin film obtained in step II. 5009


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Figure 3. (a) Schematic of the electrical transport measurement on Ag-TCNQ thin films by using EGaIn as the top-electrode material. (b) Current− voltage (I−V) characteristics of Ag-TCNQ thin films at 300 (red), 400 (blue), and 300 K (green). I−V plots were fitted linearly (yellow lines) in the voltage range of −0.3 to 0.3 V to calculate resistances RHRS and RLRS, and respective values are provided. The blue colored current (I) axis is for the 400 K plot, and the black colored current (I) axis is for both plots recorded at 300 K. (c) Plot of resistance values at HRS and LRS versus the number of cycles. (d) Heating and cooling curves of resistance versus temperature (R−T) clearly show a hysteresis loop vis-a-vis memory effect.

exchange process: (i) thickness values of initial and final thin films were markedly different (∼850 nm of Ag-TCNQ at the cost of ∼550 nm of Cu-TCNQ) (Figure S3) and (ii) the CuTCNQ thin film was completely etched away upon dipping into AgNO3 solution for 5 h (long time is a prerequisite for a classical ion-exchange process) (Figure S4). Then what could be the possible mechanism behind the sacrificial growth? To understand the origin of such an unusual transformation, we have closely examined the growth of the thin film in each cycle of the LbL by FESEM and EDXS analyses. In the first cycle, upon dipping the Cu-TCNQ thin film into AgNO3 and TCNQ solution, we observed the formation of small Ag nanoparticles (NPs) distributed over the nanoflakes of CuTCNQ (Figure S5). Over progressive cycles, Cu-TCNQ nanoflakes gradually vanished and eventually Ag-TCNQ nanorods were formed, thereby justifying the role of CuTCNQ as a sacrificial layer in the growth of Ag-TCNQ via the LbL approach. Two important reactions behind the sacrificial LbL are proposed as

(Figure 2d). EDXS analysis was further complemented by X-ray photoelectron spectroscopy (XPS) data recorded on the thin film obtained in step II (Figure 2e). The Ag 3d5/2 (3d3/2) XPS signal at a binding energy value of 373.2 (367.53) eV confirmed the presence of a Ag(I) ion, which is typical of AgTCNQ, and the Cu(I) ion was found to be absent (no reasonable Cu 2p XPS signal at the binding energy range of 930−960 eV).23 Raman spectra of both initial and final thin films showed the presence of vibrational bands at 2205, 1602, and 1374 cm−1, corresponding to CN stretching modes, CC (ring) stretching modes, and C−CN stretching modes, respectively, which are characteristics of a monoanionic TCNQ ligand (Figure S2).15 Considering the structural similarities of Cu-TCNQ and AgTCNQ, one would have expected the formation of heterostructured Ag-TCNQ/Cu-TCNQ via LPH. Surprisingly, our EDXS, XPS, Raman, and PXRD data all together suggest formation of a Ag-TCNQ thin film at the end where the initial Cu-TCNQ thin film is sacrificed over continued cycles. The here discovered Cu-TCNQ → Ag-TCNQ transformation in the thin film configuration apparently appears to be typical of an ion-exchange phenomenon and/or postsynthetic modification (metal ion metathesis) in the domain of CPs. For example, a Mn-based CP when subjected to an ionexchange process with Fe, Co, Ni, and Cu ions resulted in the formation of respective isomorphous CPs, which were otherwise difficult to obtain by standard solvothermal synthesis.24−26 Also, taking soft−hard acid−base (SHAB) principles into account, a kinetically labile metal−ligand exchange reaction is expected because both Cu(I) and Ag(I) ions are soft acids and TCNQ− is a soft base.27 However, the following observations ruled out the possibility of an ion-

CuTCNQ + AgNO3 → Cu 2 + + Ag 0 + NO3− + TCNQ− Ag 0 + TCNQ → AgTCNQ

The Cu(I) ion acted as a reducing agent for the Ag(I) ion (E0Cu(II)/Cu(I) = 0.16 V and E0Ag(I)/Ag = 0.80 V) and was sacrificed to the Cu(II) ion.28 Interestingly, we performed some control experiments where Cu(II) and Cu(I) ions were separately added to TCNQ solution, and only in the case of the Cu(I) ion we could isolate some solid Cu-TCNQ, perhaps due to the fact that the Cu(II) ion is a borderline acid (Figure S6). Thus, our sacrificial LPH not only yielded high quality but 5010


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The Journal of Physical Chemistry Letters also generated a semiconducting Ag-TCNQ thin film on the SAM template. Note that the growth of thin films of MTCNQs without the use of M as a precursor with a substrate remained challenging, specifically for Ag-TCNQ.29,30 The here presented sacrificial LbL approach appears to be the first report on successful growth of Ag-TCNQ thin films on organically modified substrates. Metal-TCNQs solids are an interesting class of air-stable semiconducting CPs with unusual magnetic properties (longrange ordering at high temperatures) as well as bistable electrical characteristics (high-impedance to low-impedance states, so-called resistive switching).31−34 So far, only optical and electrical field-induced resistive switching was observed in metal-TCNQs and other systems.35 Herein, thermally driven resistive switching in thin films of CPs is reported for the first time. Electron transport in the Ag-TCNQ thin films was initially investigated (Keithley 4200 SCS Parameter Analyzer) by using eutectic gallium indium (EGaIn) alloy as top electrodes36 (Figure 3a). The appreciable feature was nonlinear (non-Ohmic) current−voltage (I−V) characteristics of the EGaIn/Ag-TCNQ interface at room-temperature (RT; 300 K). The system was initially in a high-resistance state (HRS) and with a gradual increase in temperature from 300 to 400 K; a transition from the HRS to low-resistance state (LRS) was observed. Also, upon cooling the system back to 300 K, complete conversion of the LRS to HRS took place, that is, a clear signature of thermally driven reversible resistive switching (HRS ↔ LRS), as shown in Figure 3b. To calculate the resistances at both the HRS and LRS, data points in the respective I−V curves in low-voltage regions (−0.3 to +0.3 V) were fitted linearly (Figure 3b). The RHRS and RLRS values were estimated to be ∼106 and ∼10 Ω, respectively. A remarkable enhancement in the electrical conductanceby a record factor of ∼105seems comparatively better than earlier reports on the electric field-induced resistive switching in Ag-TCNQ.37 To test the durability of our temperature-induced resistive switching in Ag-TCNQ thin films, heating and cooling processes were performed up to 10 consecutive cycles. Excellent retention of the conductance values at both the HRS and LRS in the cycling process was observed (Figure 3c). Also, the appearance of a prominent hysteresis loop in the heating−cooling path is characteristic of the memory effect (Figure 3d). The discrete transition between the HRS and LRS can be assigned to OFF and ON states, thereby making our Ag-TCNQ thin film a promising candidate for the development of temperature-dependent memory devices. A good number of mechanisms accounting for the resistive switching in various metal oxides as well as metal−organics are proposed.31,38 In a broad sense, it is primarily either chemical (charge transfer/ conductive path via filament formation) or physical (structural/ conductive path via defect migration). As for the electrical fieldinduced resistive switching effect in the Ag-TCNQ system, formation and rupture of electronically conductive filaments (CFs) are well-established via the following charge transfer.39

Figure 4. (a) Variable-temperature out-of-plane XRD patterns of the Ag-TCNQ thin film. (b) Temperature-dependent Raman measurements on the Ag-TCNQ thin film.

subsequently cooling down to 300 K, no reasonable change in the XRD pattern was observed, thereby ruling out the structural phase transition to be in the origin of thermally driven resistive switching. In Raman spectra, the appearance of vibrational bands at ∼1374 cm−1 (C−CN stretching), ∼1603 cm−1 (CC ring stretching), and ∼2206 cm−1 (CN stretching) in our AgTCNQ thin film are characteristics of a TCNQ radical anion in the HRS (Figure 4b).15 Upon increasing the sample temperature from 300 to 400 K, the spectral fingerprints remained unchanged at the LRS, and specifically the appearance of new bands at ∼1457 and ∼2229 cm−1, characteristics of neutral TCNQ, were not observed.39 Also, with subsequent cooling down of the sample to 300 K, spectral features were retained. Thus, our Raman data clearly suggest that thermally driven resistive switching was not due to formation of CFs involving charge transfer, as discussed earlier. Notably, in a recent report on the electric field-induced resistive switching in thin films of CPs, formation of CFs was also excluded.40 Another source of the resistive switching could be the metal−semiconductor interface, the so-called interface-type.38 In order to study the role of contact resistance, we have used Ti and Pt micropads, in addition to EGaIn drops, for electrical measurements on the Ag-TCNQ thin films. Ag-TCNQ is an ntype semiconductor having an effective work function value (ϕ) of ∼1.19 eV. Because the ϕ value of Ag-TCNQ is much lower than those of EGaIn (ϕ ≈ 4.2 eV), Ti (ϕ ≈ 4.3 eV), and Pt (ϕ ≈ 6.3 eV), at each of the metal−semiconductor interfaces, a Schottky-type barrier is expected, giving rise to a depletion layer due to a band bending phenomenon (Figure 5a).30,36,38,41 Indeed, I−V plots (−5 to +5 V) recorded on various interfacial systems (EGaIn/Ag-TCNQ, Ti/Ag-TCNQ, and Pt/AgTCNQ) consistently showed non-Ohmic characteristics at 300 K (Figure 5b−d). Upon heating the systems to 400 K, nonOhmic conduction consistently switched to Ohmic conduction, and again, cooling down the systems to 300 K, non-Ohmic conduction consistently appeared back. To evalute the contact resistance values at 300 and 400 K for each of the interfacial systems, data points in the I−V curves in low-voltage regions (−0.3 to +0.3 V) were fitted linearly (Figure 5b−d, yellow lines). The respective values at 300 and

[Ag‐TCNQ]n (HRS) ↔ x Ag 0 + x TCNQ0 + [Ag‐TCNQ]n − x (LRS)

To investigate the mechanism of our thermally driven resistive switching phenomenon, we have recorded the XRD patterns and Raman spectra of the Ag-TCNQ thin film at various temperatures (Figure 4a). The XRD pattern of our AgTCNQ thin film at 300 K suggests a phase-I structure. Upon gradually increasing the sample temperature up to 400 K and 5011


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Figure 5. (a) Schematic of the energy diagram of the metal−semiconductor interface. Here, metal represents EGaIn or Pt or Ti and semiconductor represents Ag-TCNQ (n-type), where ϕm is the metal work function, ϕs is the semiconductor work function, ϕSB is the Schottky barrier, χs is the electron affinity of the semiconductor, Ec is the energy of the conduction band minimum, Ev is the energy of the valence band maximum, and Evac is the vacuum energy. (b−d) I−V characteristics of Ag-TCNQ thin films by using different top electrodes such as EGaIn (b), Ti (c), and Pt (d) at 300 (red), 400 (blue), and 300 K (green). (Inset in d: optical image of Pt contact pads on a Ag-TCNQ thin film). The blue colored current (I) axis is for the 400 K plot, and the black colored current (I) axis is for both plots recorded at 300 K (b−d).

Figure 6. (a) I−V plots of bulk Cu-TCNQ (in pellet form) recorded at different temperatures, 300 (red), 370 (blue), and 300 K (green), with direct contacts of pins, the so-called two-probe. (b) I−V plots of the Cu-TCNQ pellet recorded at different temperatures, 300 (red), 370 (blue), and 300 K (green), with EGaIn as the top electrodes. The blue colored current (I) axis is for the 400 K plot, and the black colored current (I) axis is for both plots recorded at 300 K (a,b). (c) C−V curves recorded on a Cu-TCNQ pellet at different temperatures, 300 (black), 370 (blue), and 300 K (red), with EGaIn as the top electrodes. (d) C−V curves recorded on Ag-TCNQ thin films at different temperatures, 300 (black), 370 (blue), and 300 K (red), with EGaIn as top electrodes.

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energy. Specifically, thin films of Ag-TCNQ and Cu-TCNQ consistently showed reversible switching between the HRS at 300 K and LRS at elevated temperatures (400 K for the former and 370 K for the later). A remarkable enhancement factor in the electrical conductance with an excellent switching durability is observed. The on-surface transformation of Cu-TCNQ → Ag-TCNQ was realized to be a sacrifical growth and not an ionexchange process. Our findings not only provide new insights into the interfacial growth of thin films of CPs but also direct toward engineering metal−semiconductor interfaces, which could bring novel electronic and magnetic properties upon physical perturbation.

400 K were found to be ∼106 and ∼10 Ω (for EGaIn contacts), ∼ 107 and ∼10 Ω (Ti contacts), and ∼105 and ∼10 Ω (for Pt contacts). Depending on the metallic contact, the enhancement factor in electrical resistance ranged from ∼104 to ∼106. Usually, in the case of metal/n-type semiconductor interfaces, the contact resistance is expected to increases with increasing work function of the metal.38 However, in the present study, no correlation among contact resistance values and the work function values of metals was observed, which strongly suggests that the Fermi levels of various metals on Ag-TCNQ are pinned, possibly due to creation of metal-induced gap states.42,43 Such a thermally driven resistive switching effect is not specific for the Ag-TCNQ system; the Cu-TCNQ thin films grown by the LbL approach on SAM templates also showed the effect (Figure S7). No appreciable change in the temperaturedependent XRD patterns and Raman data on the Cu-TCNQ thin films consistently excluded the formation of CFs as the driving mechanism for the thermally driven resistive switching (Figure S7). To further strengthen our claim of thermally driven resistive switching as an interfacial effect, I−V plots on a pressed pellet of powder Cu-TCNQ were recorded at various temperatures using direct contacts of the pins (without the use of EGaIn as top contacts) and also with the help of EGaIn as top contacts (Figure 6a,b). I−V features in Figure 6a,b were remarkably different: (i) in the former system, Ohmic characteristics were at both 300 and 370 K, whereas in the later system, non-Ohmic (signature of Schottky barrier) and Ohmic characteristics were at 300 and 370 K, respectively; (ii) the contact resistance value across the EGaIn/Cu-TCNQ interface at 300 K is significantly higher (by a factor of ∼103) than the direct contact resistance; and (iii) a clear onset of thermally driven resistive switching only at the EGaIn/CuTCNQ interface with an enhancement factor of ∼103 in the electrical resistance value. Thus, the observed thermally driven resistive switching mainly originates from the lowering of the Schottky barrier at the metal/semiconductor interface by increasing the temperature44,45 and not a bulk property of the material itself. The Schottky barrier height can be approximated with the following equation:45 ϕSB(T) = ϕSB(300 K) + α(T − 300 K), where α is a temperature coefficient term (eV/K) and for n-type semiconductors such as Ag-TCNQ and Cu-TCNQ, the values of α are negative. Finally, we have estimated potential profiles of the depletion layer by recording capacitance−voltage (C−V) plots at 300 and 370 K (PARSTAT MC PMC-2000). In a conventional Schottky model, C is inversely proportional to the depletion layer width, and thus, C is expected to be larger in the LRS than that in the HRS.38 Indeed, our preliminary C−V data recorded on both EGaIn/Cu-TCNQ (pellet) and EGaIn/Ag-TCNQ (thin film) interfaces showed larger C values at hightemperatures (LRS) in comparison to C values at roomtemperature (HRS) (Figure 6c,d). However, to conclude on the temperature dependency of the Schottky barrier height from C−V measurements, more analysis is required. We have also studied the time dependence of thermally driven resistive switching in Ag-TCNQ thin films. The system was heated to 400 K, the resistance value was monitored at every 60 s time interval while gradually cooling down to 300 K, and the memory time of the LRS was realized to be 720 s37 (Figure S8). In conclusions, we have successfully demonstrated resistive switching in thin films of CPs upon alternation of the Schottky barrier at the metal−semiconductor interfaces by thermal

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ASSOCIATED CONTENT

S Supporting Information *

The Supporting Information is available free of charge on the ACS Publications website at DOI: 10.1021/acs.jpclett.7b02138. Experimental section with characterizations and supplementary figures (PDF)

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AUTHOR INFORMATION

Corresponding Author

*E-mail: [email protected]. ORCID

Nirmalya Ballav: 0000-0002-7916-7334 Notes

The authors declare no competing financial interest.

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ACKNOWLEDGMENTS Financial support from IISER Pune, SERB (EMR/2016/ 001404), and MHRD-FAST (India, Project − CORESUM) is thankfully acknowledged. S.R. and P.K.J. thank IISER Pune for providing Senior Research Fellowships. A.P. thanks DST-KVPY for a fellowship. The authors thank Dr. G. V. Pavan Kumar for Raman data; Mr. Parveen Nasa, Mr. Nilesh Dumbre, and Dr. S. Singh for XRD measurements; and Dr. Shouvik Datta and Prashant Kale for photolithography and sputtering. The authors sincerely acknowledge the UGC-DAE Consortium for Scientific Research, Indore (India) for providing variable-temperature Raman data and thank Dr. Vasant Sathe for the support.

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REFERENCES

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Thermally-Driven Resistive Switching in Solution Processable Thin

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