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Nanophases in mechanochemically synthesized AgI–CuI system: structure, phase stability and phase transitions
Journal of Physics and Chemistry of Solids 65 (2004) 1669–1677 www.elsevier.com/locate/jpcs
Nanophases in mechanochemically synthesized AgI– CuI system: structure, phase stability and phase transitions D. Bharathi Mohan, C.S. Sunandana* School of Physics, University of Hyderabad, Hyderabad 500 046, India Received 20 October 2003; revised 24 March 2004; accepted 9 April 2004
Abstract Nanoscale crystallites of Ag-rich (Ag12xCuxI, x ¼ 0:05; 0.10, 0.15 and 0.25), Cu-rich (Cu1-yAgyI, y ¼ 0:05; 0.10, 0.15 and 0.25) and intermediate Ag1-xCuxI ðx ¼ 0:50Þ solid solutions and end members AgI, CuI with sizes in the range of 46 –13 nm were synthesized by attrition at ambient temperature in a soft mechanochemical reaction (MCR) of Ag, Cu and I. Monophasic g-AgI (zincblende, a ¼ 638 pm) with disordered Agþ sublattice and the crystallite size of about , 31 nm was realized in the case of Ag0.75Cu0.25I ðx ¼ 0:25Þ composition. Lattice parameter decreases linearly from 649 to 604 pm with increasing Cu concentration in the AgI– CuI system validating Vegard’s law. Smallest size (,13 nm) agglomerated nanocrystals were realized in the Cu-rich composition Cu0.75Ag0.25I (a ¼ 615 pm), while unagglomerated uniform-sized (,17 nm) and spherical shape nanocrystallites of Ag0.50Cu0.50I (a ¼ 626 pm) with maximum strain were synthesized for sensor applications using MCR. Differential scanning calorimetry study shows the systematic changes in the phase transition temperature with Cu substitution. Ag-rich composition posses less enthalpy (DH (x or Cu ¼ 0.05, 0.10, 0.15, 0.25) ¼ 6.0, 6.11, 6.6, 6.3 in kJ/mol) and entropy (DS (y or Ag ¼ 0.05, 0.10, 0.15, 0.25) ¼ 14.15, 14.1, 15.03, 13.6 in J/mol K) when compared to undoped AgI (DH ¼ 9:63 kJ=mol; DS ¼ 22:8 J=mol K) implying greater thermal stability of g-phase due to Cu-strengthened Ag– I bond. Enhanced entropy (DS ¼ 8:17 J=mol K) in Cu0.75Ag0.25I (Cu-rich) solid solutions relative to CuI (DS ¼ 1:0 J=mol K) indicates Ag-induced cation disorder. Fifteen percent Ag-doped CuI (Cu0.85Ag0.15I) nanocrystals apparently behave like microscopic p – n junctions with currents in the range of 1026 –1028 A characterized by a non-linear I – V curve. q 2004 Elsevier Ltd. All rights reserved. Keywords: D. Phase transitions
1. Introduction Mechanochemical reaction (MCR) or mechanical grinding has recently emerged as a viable and convenient method to produce nanoscale ceramics [1]. MCR triggered by the application of mechanical energy, is characterized by a large negative fractional energy change and is thermodynamically feasible at ambient temperature. The over all reaction can be controlled by the mechanical generation of clean/fresh surfaces due to fracturing, increased defect density and reduction in particle size. Boldyrev [2] described the main methods of investigation of the mechanochemical process in inorganic solids, as well as the influence of preliminary mechanical processing on the reactivity of * Corresponding author. Tel.: þ 91-40-23134324; fax: þ 91-4023010227. E-mail address: [email protected] (C.S. Sunandana). 0022-3697/$ - see front matter q 2004 Elsevier Ltd. All rights reserved. doi:10.1016/j.jpcs.2004.04.007
solids. Grinding generally influences texture and structure leading to a decrease of the particle size of the product phase and, simultaneously, to an increase of the microstrain due to the contribution of the grain boundaries formed during the process [3]. AgI (CuBr and CuI) [5,6] are well-studied superionic conducting materials with zincblende structure at room temperature, and rather high Agþ conductivity [4]. The p –d hybridized [7] Ag – I [8] and Cu – I [9] bonds are weakly tetrahedral, with ease of dislocations/stacking faults production in the rather soft sphalerite or zincblende lattice, with cation disorder. This unique feature is conducive for realizing metastable and new superionic phases by mechanochemistry. CuI is a mixed ionic– electronic conductor at room temperature with predominant electron –hole conduction up to 200 8C. (1) The cubic (superionic) a-AgI forms above 146.5 8C, in which iodine ions form a bcc lattice and two silver ions are distributed randomly among the many available crystallographic sites. (2) b-hexagonal
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(non-superionic) phase (Wurtzite structure) of AgI forms reversibly upon cooling the high temperature cubic form below 146.5 8C in which each iodine ion being surrounded by a regular tetrahedron of four silver ions and vice versa. (3) g-AgI, with a disordered Agþ sublattice in its zincblende structure and possessing an Agþ conductivity higher than that of b-AgI, is metastable below 137 8C [4,10]. An ambient solid-state process such as the MCR, when applied to these ‘soft’ superionic conductors AgI and CuI, can be used to synthesize new materials in terms of both composition and microstructure applicable in ‘as prepared’ form in electrochemical sensor and battery systems at ambient. In this paper, the influence of mechanical work on the constituents of Ag-rich (Ag12xCuxI) and Cu-rich (Cu12yAgyI) solid solutions and the end members AgI, CuI has been systematically investigated in a MCR using XRD, scanning electron microscope (SEM), differential scanning calorimetry (DSC) and electronic conductivity as probes. A new phase (Ag0.50Cu0.50I) has been obtained using this unique but convenient technique. A variety of elementary physicochemical processes at micro- and macro-levels are expected to be stimulated leading to such metastable nanocrystalline products with optimal (micro) structure and morphology.
2. Experimental methods Compositions corresponding to AgI, Ag-rich (Ag12xCuxI where x ¼ 0:05; i.e. 5%, 0.10, 0.15, 0.25), intermediate Ag12xCuxI ðx ¼ 0:50Þ; Cu-rich (Cu12yAgyI, y ¼ 0:05; 0.10, 0.15, 0.25) portion of the AgI –CuI system and CuI were synthesized by mechanical grinding in a 600 agate mortar and pestle for 5 h at room temperature in an unilluminated room by using appropriate quantities (in wt%) of copper (LOBA, India), silver (Special materials project, Hyderabad, India) and I (Rasayana Laboratory, India). When attrition was carried out, the charge (Ag or Cu or I or Ag þ Cu or Cu þ Ag or Ag þ Cu þ I or Cu þ Ag þ I) present at the bottom of the mortar is sheared by the pestle at the start of the spiral type motion executed by it. This motion causes the charges to spread over the inner surface of the mortar; say 1/3 of the inner surface area (294.5 cm3). In about a minute of attrition the charge undergoes ‘reactive mixing’ as the pestle circles about 60 times in the mortar—a motion that is fairly uniform and periodic. The periodic motion of the pestle causes Ag/Cu grains to get compressed and the compressed surface is exposed to iodine molcules on the breaking-up and sublimating iodine flakes. Structural characterization on these samples was performed using a PHILIPS X-Ray diffractometer with Cu Ka1 radiation. Microstructure studies were carried out using PHILLIPS XL-30 SERIES SEM by spreading the assynthesized powder over a well-cleaned microscopic slide after applying vacuum grease. Phase stability and phase transitions were investigated using a DSC instrument (TA, USA) with crimped samples held in an argon atmosphere.
The powdered sample was compressed at pressures of 4 ton in a stainless steel dye at room temperature. The pellets thus obtained had an area of 8 mm and thickness of 3 mm. Pellets were annealed at 373 K(^ 1 K) for 3 h and used for measuring I – V characteristics. I – V measurements were made by a two-probe method. In order to measure the electronic conductivity, a graphite electrode was used to block the ion movement on the other side of the sample configuration. The electrode configuration in this case is M/MX/B where M ¼ Ag or Cu, MX ¼ (Ag, Cu)I, B ¼ graphite. A constant voltage source was connected in series with the sample and a fixed (standard) resistance of 10 ohm. A KEITHLEY 195A digital multimeter was used for measuring the current flow through the sample.
3. Results and discussion 3.1. Structure of b- and g-AgI and CuI, Vegard’s law and particle size Fig. 1 (expanded spectrum includes only major lines) shows the XRD patterns of as prepared AgI –CuI solid solutions. Undoped AgI sample (pattern (a)) contains two phases: the major g-AgI (sphalerite or zincblende structure, approx. 85.55%) phase, characterized by the XRD peaks (111), (220), (311), (400), (331), (422) and (511) [a ¼ 649 pm], and the minor b-AgI (Wurtzite structure, appox. 14.45%) identified by (100), (101), (102) and (103) reflections [a ¼ 412 pm; c ¼ 730 pm]. The main purpose of Cu doping is to stabilize the metastable zincblende structure
Fig. 1. X-ray diffractograms of AgI–CuI solid solutions mechanically ground for 300 min. (a) AgI; (b) Ag0.95Cu0.05I; (c) Ag0.90Cu0.10I; (d) Ag0.85Cu0.15I; (e) Ag0.75Cu0.25I; (f) Ag0.50Cu0.50I; (g) Cu0.75Ag0.25I; (h) Cu0.85Ag0.15I; (i) Cu0.90Ag0.10I; (j) Cu0.95Ag0.05I; and (k) CuI.
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of g-AgI besides assisting in particle size reduction. g-AgI nanoparticles were clearly seen upon progressively iodized Ag – Cu metastable alloy thin films under ambient conditions, which have been studied, by optical absorption and photoluminescence [11]. The Ag0.95Cu0.05I ðx ¼ 0:05Þ sample (pattern (b)) has two b-AgI lines (100) and (104) [approx. 3.6% in relative intensity] and one silver line (111) with six g-AgI lines (approx. 96.4%). Ag0.90Cu0.10I ðx ¼ 0:10Þ composition has only one b-AgI line (100) with negligibly small intensity (approx. 1.8%) [pattern(c)]. The Bragg reflections of Ag0.85Cu0.15I sample [pattern (d)] show predominant peaks corresponding to the zincblende structure of g-AgI alone. More significantly, the growth of wurtzite phase of b-AgI is now completely suppressed, which suggests a rapid nucleation of g-AgI crystallites aided by Cu. Moreover, it contains unreacted Ag (111), signaling the formation of monophasic g-AgI. But in the case of Ag0.75Cu0.25I [pattern (e)], there is no Ag line implying complete consumption of Ag, so that the pattern corresponds to monophasic g-AgI with sharper peaks. Fig. 2 shows a gradual increase of all the Bragg angles accompanied by a broadening of the (111) plane conspicuously observed upon progressive Cu addition. The lattice parameter (a in pm) decreases linearly with increasing Cu (wt%) in a continuous manner (Fig. 3). As more and more Cu atoms substitute Ag in AgI, the Ag – I bond progressively shortens, there by stabilizing the g-AgI structure. In the Cu-rich region, Ag is accumulated in the already stable zincblende CuI lattice. The systematic development and stabilization of g-CuI using grinding time as a parameter have already been discussed [12]. Undoped CuI [pattern (k)] shows clearly, the more fully developed Bragg reflections of g-CuI lines
Fig. 2. Gradual increase of Bragg angle accompanied by a broadening of the (111) plane upon progressive Cu addition. (a) AgI; (b) Ag0.95Cu0.05I; (c) Ag0.90Cu0.10I; (d) Ag0.85Cu0.15I; (e) Ag0.75Cu0.25I; (f) Ag0.50Cu0.50I; (g) Cu0.75Ag0.25I; (h) Cu0.85Ag0.15I; (i) Cu0.90Ag0.10I; (j) Cu0.95Ag0.05I; (k) CuI; and (l) commercial powder of CuI.
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Fig. 3. Plot of lattice parameter ðaÞ vs Cu composition in the AgI– CuI system of solid solutions showing smooth linear decrease of ‘a’ from AgI to CuI.
characterized by (111), (200), (220), (311), (222), (400), (331), (420) and (422) with intensities larger than those of Cu-rich (Cu12yAgyI) phases. A continuous shift in the Bragg angle is observed in all Bragg peaks [patterns (j), (i), (h) and (g)] of Cu-rich system as it moves towards Ag-rich region (Fig. 1) upon progressive addition of Ag. From Fig. 3, we observe a monotonic linear increase in the lattice parameter (a in pm) of all compositions relative to undoped CuI. The smaller ionic size of Cuþ (145 pm) as compared to Agþ (165 pm) probably results in the systematic increase in the lattice parameter (a) in a linear fashion with increasing Ag addition [13,14]. The Bragg reflections of the intermediate (50:50) composition Ag0.50Cu0.50I (where x ¼ 0:50) lie in-between Ag0.75Cu0.25I (pattern (e)) and Cu0.75Ag0.25I (pattern (g)). The changes in 2u of the first three prominent peaks are, 24.38 (111); 40.41 (220); 47.77 (311). Also the intensities of (200), (222) and (420) lines become negligibly small. The smooth decrease in ‘a’ (pm) (or unit cell volume) of AgI is linear as Cu concentration is increased (Fig. 3). This indicates a continuously uniform stabilization of static cation disorder in the sphalerite or zincblende AgI lattice brought about by the mechanochemical processing. Finally, Vegard’s law seems to be valid in AgI – CuI system [15]. All samples produce appreciable diffraction broadening and it is reasonably assumed that this arises from small crystallites and internal stresses. Diffraction theory predicts that the breadth due to crystallite size varies with angle as sec u and that due to elastic strain as tan u [15]. The additional broadening in diffraction peaks beyond the inherent peak widths due to instrumental effects can be used to measure crystallite sizes as low as 1.0 nm. The crystallite sizes of as-prepared compositions were calculated from the full width at half maximum (FWHM) of three prominent peaks (111), (220) and (311) using the Debye –Scherrer formula, t ¼ ½ð0:9lÞ=ðB cos uÞ [16].
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Use of the Lorentzian function to fit the FWHM of three peaks yields more accurate (error is approx. 0.22%) results than those obtained by using Gaussian (1%) and Voigt functions (0.5%). Therefore, the observed XRD lines are essentially Lorentzian, pointing to a uniform crystallite size distribution. These apparent sizes essentially correspond to the true size of the crystallites, because the observed induced microstrain is generally very small (, 1023 – 1024) due to the relatively small mechanical energy input to the system during the MCR process. Furthermore, the elastic constants (particularly the shear elastic constant (C44)) of Ag (0.46 £ 1012 dynes/cm2) and Cu (0.75 £ 1012 dynes/ cm2) at 300 K [17] are very close and are relatively small, low power, soft attrition process does not induce much strain in the fully reacted solid solution samples. Overall, we can say that the broadening is due to the microstrain presence in the crystallites is negligible. Effects of Cu addition in reducing the particle size in the Ag12xCuxI solid solutions prepared by aqueous solution method in our laboratory were discussed previously [14]. Fig. 2 shows the gradual increase of Bragg angle accompanied by a broadening of the (111) plane upon progressive Cu addition in the Ag-rich region. Fig. 4 clearly shows the crystallite size vs Cu composition plot for the AgI –CuI system. Bigger size (, 46 nm) crystallites were observed in undoped AgI, because it contains both b-AgI (, 14.5%) and g-AgI (, 85.5%) phases. 5% Cu addition shows a sudden fall in the crystallite size (, 31 nm). The smaller Cuþ ions are expected to be present mostly in the subsurface region of the crystals rather than in the crystal bulk, facilitating further reduction in the interfacial energies between the randomly distributed polytypes of AgI, leading to the stabilization of the metastable cubic zincblende g-AgI with smaller crystallites [18]. Interestingly, the smallest crystallite size namely,
Fig. 4. Crystallite size changes in AgI–CuI solid solutions made by mechanochemical reaction. Smallest crystallite size occurs for Cu0.75Ag0.25I but unagglomerated crystallites of 17 nm are seen in Ag0.50Cu0.50I.
26 nm is observed in the Ag-rich region, which corresponds to the Ag0.90Cu0.10I solid solution. It is significant that Ag0.75Cu0.25I ðx ¼ 0:25Þ solid solution contains slightly bigger crystallites of about 31 nm, due to the absence of strain that is normally induced by the MCR process. It is interesting to look at the strain analysis plot where x ¼ 0:25 of Cu-doped system shows a horizontal straight line (see later) in contrast to undoped AgI and Ag-rich (x ¼ 0:05 and 0.10 Cu doped) compositions. In the case of Cu-rich samples, increase in Ag addition leads to a systematic broadening of the XRD peaks (Fig. 2 shows the (111) plane). So the particle size also starts decreasing from undoped CuI (, 31 nm). The Cu-rich composition Cu0.75Ag0.25I ðy ¼ 0:25Þ shows the broadest peak, corresponding to the smallest particle size (, 13 nm) observed, due to the presence of strain, stacking faults and grain boundaries. This implies that shear strain and defects together reduce the size of the nanoparticles synthesized by the MCR process. It is interesting that even 5% Cu substituted into AgI brings about a considerable reduction in the particle size. This can explained by noting that the polarizing power of the smaller-sized Cuþ is greater than that of Agþ [19]. Also the Cu – I bond is more covalent than the Ag – I bond. For these two reasons, the presence of even a small amount ($ 5%) of Cu in AgI initiates the particle size reduction processes, one of which may be the considerable reduction of forces responsible for agglomeration. So, regardless of whether one is dealing with Cu-rich or Ag-rich compositions, the particle size reduction invariably seems to take place. 3.2. Strain analysis The Nelson – Reily function (NRF) vs Dd=d curves [20,21] and least squares regression line: ð2wÞ2 cos2 u as a function of sin2 u [22] for AgI –CuI system obtained from the XRD data are plotted in Fig. 5a,b, respectively. Systematic lattice transformation of all crystal planes upon Cu substitution is clearly evident (Fig. 5a). Dd=d values increase up to x ¼ 0:50 and thereafter it starts decreasing all the way down to undoped CuI ðy ¼ 0Þ: In the case of Ag-rich, Cu-rich (y ¼ 0:05; 0.10) and undoped CuI samples, Dd=d values lie on a horizontal straight line, almost parallel to NRF axis, which indicates the absence of stacking faults in the cubic structure. But, rather large scatter observed in Dd=d values for the intermediate 50– 50 (Ag0.50Cu0.50I) and Cu-rich (Cu0.75Ag0.25I, Cu0.85Ag0.15I) samples is a measure of the high concentration of stacking faults present in these compositions. Ag-rich and Cu-rich samples show less scatter with a small slope value whereas intermediate (50, 75% Cu doped) compositions gives large slope value (Fig. 5b). Average microstrain is calculated and tabulated for the all compositions by measuring slopes of the least squares regression line (Table 1). Maximum strain values are observed only in the intermediate 50 and 75% Cu
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Fig. 5. (a) Strain analysis of AgI–CuI solid solutions obtained from XRD. Strain analysis of AgI–CuI solid solutions obtained from XRD. (A) AgI; (B) Ag0.90Cu0.10I; (C) Ag0.75Cu0.25I; (D) Ag0.50Cu0.50I; (E) Cu0.75Ag0.25I; (F) Cu0.85Ag0.15I; (G) Cu0.95Ag0.05I; see text for details. (b) Least square fit analysis shows the induced microstrain obtained from XRD. Slope gives the average strain value. (a) Ag0.75Cu0.25I; (b) Ag0.50Cu0.50I; (c) Ag0.10Cu0.90I; and (d) Ag0.85Cu0.15I. Other compositions (not marked) do not show significant strain.
doped AgI compositions but not in others. The larger covalency of CuI (relative to AgI) [23] possibly assists in generating considerable strain in Cu-rich compositions leading to size-reduced and unagglomerated Ag0.50Cu0.50I crystallites as seen in SEM. This is in conformity with the results of particle size analysis where, these compositions have minimum particle size. Less scattered points are observed in pure AgI system due to the presence of hexagonal (b-AgI) structure. There are two stages of structural transformation observed in the presence of stacking faults in this system one being the transformation of the mixture of minor Wurtzite (b-AgI) and major zincblende (g-AgI) phase to pure zincblende (g-AgI) phase and the other being the pure g-AgI phase to zincblende (g-CuI) phase transition. The latter one is essentially a weakly covalent to a strongly covalent bonding transition taking place gradually in the Cu-rich phases. 3.3. Uniform size MCR Ag0.50Cu0.50I crystallites (13 nm) Secondary electron images show qualitatively the differences in the surface morphology of mechanochemically reacted ground samples of AgI, Ag0.75Cu0.25I, Ag0.50Cu0.50I, Cu0.75Ag0.25I, Cu 0.90 Ag 0.10 I and CuI (Fig. 6). Random orientation of bigger size (, 46 nm) crystallites were observed in 300 min undoped AgI sample (Fig. 6a). As the concentration of doping element (25%Cu) is increased, the crystallite size also decreases to , 31 nm, so that for middle composition (Ag0.50Cu0.50I), crystallites with nearly uniform size (, 17 nm) (Fig. 6c) in a completely formed solid solution were observed indicating maximized
contact areas between Ag and Cu leading to optimal reactivity. The formation of this new composition Ag0.50Cu0.50I is of both fundamental and technical interest because of the difficulty of using conventional techniques such as fusion and solid-state reaction for its realization. Agglomerated crystallites with small size (, 13 nm) and increased reactive surface area were observed in the (Cu-rich) Cu0.75Ag0.25I composition (d). Unreacted Ag (bigger particles in Fig. 6e) with agglomerated Cu particles was observed in Cu0.90Ag0.10I ðy ¼ 0:10Þ sample. Eventhough, the MCR technique is a non-equilibriumprocessing route, our SEM (and XRD) observation suggests that apart from incorporating residual strain, the samples obtained are by and large homogeneous. The relatively low energy cost for, and the high efficiency of, the process could lead to homogeneous samples. Table 1 Average apparent crystallite size and strain Compositions
Crystallite size (nm)
Microstrain (1024)
AgI Ag0.95Cu0.05I Ag0.90Cu0.10I Ag0.85Cu0.15I Ag0.75Cu0.25I Ag0.50Cu0.50I Ag0.25Cu0.75I Ag0.15Cu0.85I Ag0.10Cu0.90I Ag0.05Cu0.95I CuI
45.7 31.0 25.6 27.0 31.4 17.0 13.0 17.1 24.0 26.7 30.6
– – – 3.05 – 3.15 18.35 – 3.76 – –
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ðDSÞ—at the phase transitions are given by DH ¼ kA=m where k is the calibration constant, A is the area under the curve and m is the mass of the sample DS ¼ DH=Tt
Fig. 6. Secondary electron micrographs for MCR derived compositions in the AgI–CuI system. (a) AgI; (b) Ag0.75Cu0.25I; (c) Ag0.50Cu0.50I; (d) Cu0.75Ag0.25I; (e) Cu0.90Ag0.10I; (f) CuI. Note: uniform sized, unagglomerated nanocrystals in (c) and smallest, agglomerated nanocrystals in (d).
3.4. Thermal analysis, enthalpy and entropy changes Fig. 7 shows the thermal response of MCR derived AgI – CuI solid solutions scanned at a rate of 10 K per minute in the non-isothermal mode. The changes in thermodynamic parameters—enthalpy ðDHÞ and entropy
Fig. 7. AgI–CuI system shows the systematic variation of phase transition temperature ðTt Þ curve, where (A) AgI; (B) Ag 0.95 Cu 0.05 I; (C) Ag0.90Cu0.10I; (D) Ag0.85Cu0.15I; (E) Ag0.75Cu0.25I; (F) Ag0.50Cu0.50I; (G) Cu0.75Ag0.25I; (H) Cu0.85Ag0.15I; (I) Cu0.90Ag0.10I; (J) Cu0.95Ag0.05I; and (K) CuI.
where Tt is the transition temperature. On heating undoped AgI, a sharp and intense endothermic peak, g ! a phase transition ðTt Þ observed at 422.6 K (420 K is in commercial AgI). The increased Tt probably arises from the relatively strengthened Ag – I bond, an effect essentially induced by the MCR process. In Ag-rich region, Ag –I bond is further strengthened by the Cu substitution. So that MCR of (Ag– Cu) I extends the transition temperature ðTt Þ range over which the g-phase is stable (Fig. 8). This extended thermal stability of the g-phase is important in view of the potential applications. A competition between Cu-induced surface area enhancement and thermally induced surface area reduction may lead to an increased cation-sublattice melting temperature reflected in DSC curves. AgI is cation disordered even at room temperature and also a weakly covalent bonded structure. The addition of Cuþ has perhaps resulted in a considerable decrease in the number of Frenkel defects with the maximum strain field abruptly produced at the g – a phase transition leading to a broad and increased asymmetric endothermic peak that systematically develops in x ¼ 0:10; 0.15, 0.25 (Ag-rich) solid solutions. It is the interaction of the interstitial cation disorder with the strain field associated with this net volume change which is the principal factor for the Cu concentration-dependence of the phase transition temperature ðTt Þ [24]. Thus, grinding causes local changes in temperature and free energy that are tied up with the crystallite size reduction. ‘Melting’ of the sublattice of mobile ions generally occurs at the temperature of the superionic
Fig. 8. AgI–CuI system shows the changes in the phase transition temperature ðTt Þ. Whereas pure AgI (at 0 wt%), Ag-rich and Cu-rich compositions exhibit g ! a transition and pure CuI (at 100 wt%) exhibits both g ! b and b ! a transitions.
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phase transition when some of the SIC ions are suddenly found in intersitial sites (Frenkel defects), where they are distributed over many vacancies (order –disorder phase transition). This sudden increase in the Frenkel defect concentration contributes significantly to the total entropy of the superionic transition. The entropy of the crystallization of materials such as Cu, Ag, Au, and Li in a ‘molten’ sublattice is explained by Korzhuev [25]. Fig. 9 shows the enthalpy/entropy vs Cu substitution in weight % at the phase transition temperature ðTt Þ: Undoped AgI shows the large value of enthalpy and entropy (DH ¼ 9:63 kJ=mol; DS ¼ 22:8 J=mol K) due to the presence of large number of Frenkel defects induced by the MCR process. Even a 5% addition of Cu causes a sudden fall (followed by a more gradual decrease) in the values of enthalpy (DH (x or Cu ¼ 0.05, 0.10, 0.15, 0.25) ¼ 6.0, 6.11, 6.6, 6.3 in kJ/mol) and entropy (DS (x or Ag ¼ 0.05, 0.10, 0.15, 0.25) ¼ 14.15, 14.1, 15.03, 13.6 in J/mol K) observed. These observations imply that the Cuþ ion ‘arrests’ the order –disorder phase transition in AgI, which means that, the number of Frenkel defects is decreased in the Ag-rich solid solutions. For the middle Ag0.50Cu0.50I composition, a single-phase transition is observed at 548.6 K, which is to be compared to 420 and 648 K (g ! b transition). The Cu –I bond is more covalent than the Ag –I bond and the strength of the former is 55 meV compared to 36 meV for the latter. On heating 5-h ground undoped CuI, two endothermic peaks are observed, one at 648 K corresponding to g ! b phase transition and the other at 671.7 K representing the b ! a phase transition of CuI. Increased Tt (DT < 6 K) in the g ! b phase transition and decreased Tt (DT < 8 K) in the b ! a phase transition are observed when compared to the phase transitions in the commercial powder CuI. This suggests that while the MCR process strengthens the g-CuI (zincblende, F4¯3m) structure at the same time it weakens the a-CuI (zincblende, Fm3¯m)
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structure. Attrition essentially prepares the fcc Cu lattice so that it favors the accommodation of I ions in a weakly covalent arrangement [26]. The undoped CuI shows very less enthalpy (DH ¼ 0:5 kJ=mol) and entropy (DS ¼ 1:0 J=mol K) compared to Cu-rich iodide solid solutions (see below). As 5% Ag is added to CuI, only one peak is observed at 649.7 K, which corresponds to g ! a phase transition implying that the g ! b transition is completely suppressed. The occurrence of a single-phase transition in CuI probably means that the difference between the structures of the b-phase and the g-phase has had no significant effect on the nature of the phase transition. As the Cu concentration is increased, a continuous decrease in the phase transition temperature was observed (Fig. 8). Also, there is no trace of the g ! a characteristic of AgI phase transition. Therefore, all Agþ ions have apparently been incorporated into the CuI lattice, partially replacing Cuþ ions in the cation-sublattice. The effects of Ag in Cu-rich solid solutions are to cause an increase in the surface energy of the nanoparticle and to induce an increase in the cation disorder in the Cuþ-ion sublattice. Consequently, larger values of enthalpy (DH (y or Ag ¼ 0.05, 0.15, 0.25) ¼ 5.8, 6.23, 4.917 in kJ/mol) and entropy (DS (y or Ag ¼ 0.05, 0.15, 0.25) ¼ 8.9, 9.95, 8.17 in J/mol K) are observed in Cu-rich solid solutions (Fig. 9). Phase transitions are not multiply reversible implying ‘thermal annealing’ of strains induced by the MCR process. The kinetics of phase transitions and the structural relaxation of MCR samples of AgI –CuI solid solutions are interesting problems that deserve to be studied separately in depth. 3.5. I – V curve The process of MCR introduces a number of steps in the formation of solid solutions; the important one being charged interface formation. These interfaces probably act as microscopic p – n junctions, and the process eventually produce a non-negligible concentration of electronic current carriers in the MCR derived nanocrystalline powder. Cu-rich compositions tend to contain holes and Ag-rich compositions contain electrons as current carriers in the I–VII mixed semiconductors (while AgI is an n-type semiconductor, CuI is a p-type semiconductor). To check if MCR samples indeed possess non-linear I – V characteristics, a dc two-probe Hebb– Wagner [27,28] polarization cell of the type, ð2Þ AglAg12x Cux IlC ðþÞ; ð2Þ CulCu12x Agx IlC ðþÞ;
Fig. 9. AgI–CuI system shows the variation of Enthalpy (V) and entropy (†) in both heating (endotherm) and cooling (exotherm) cycles.
was constructed. When a dc potential (V) less than the decomposition potential of the electrolyte is applied with positive polarity on the right side, the Agþ ions move initially from the right to Ag electrode. Graphite acts as a blocking electrode so that there is no supply of Agþ ions at the ClAgrich interface. At the steady-state, the potential gradient due to electric field is just balanced by that due to the chemical
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potential gradient so that the ionic current is zero. Under these conditions, the current is solely due to differential flow of electrons or holes. Wagner has shown that the steady-state current (I) is given by I ¼ Ie þ Ih ¼ ðkTA=eLÞ½se {1 2 expð2eV=kTÞ} þ sh {expðeV=kTÞ 2 1} where se ; sh are the conductivity due to electrons and holes, respectively. L is the length, A the area of the sample, k; the Boltzman constant and V is the potential. Ag-rich solid solutions show the systematic variation in the I – V curve (Fig. 10a). Pure AgI shows the linear curve with the trivial current flow in the order of 1029 –1028 A. There are no significant effects observed up to 10% Cu substitution. The Ag0.85Cu0.15I solid solution shows a ramp-like current behavior up to 0.35 V and nonlinearity just shows up thereafter. The sudden rise in the current observed after 0.35 V in the order of 1028 –1027 A in a non-linear manner in the Ag0.75Cu0.25I solid solution,
Fig. 10. (a) I – V characteristics of pure AgI and Ag-rich solid solutions at 300 K. (b) I – V characteristics of intermediate 50–50 solid solution at 300 K.
suggests that a substantial number of holes are injected by Cu in the Ag-rich solid solutions. The current behavior was found to be different for the two electrodes Ag and Cu used in the experiments. Fig. 10b shows the non-linear response of current with increasing voltage for the Ag0.50Cu0.50I composition. Though, there is no noticeable change in the current behavior up to V < 0:30 V; there is a significant difference between the characteristic V – I behavior of the sample with Ag as electrode and with Cu as electrode after V $ 0:35 V: While, it is the residual current behavior below V < 0:30 V; the change in current behavior at the threshold voltage of 0.35 V implies that the voltage-induced hole current response similar to that of a forward biased p –n junction. Fig. 11 shows the electronic conductivity for pure CuI and Cu-rich solid solutions. On the other hand, CuI (commercial powder) gives a nearly linear curve with the current changing from 1026 to 1024 A. But there are two current steps found in the MCR sample CuI: one at < 0.2 V and the other at 0.33 V. Thereafter, it is less linear than I – V plot of the commercial powder. The obtained current magnitude is one order higher than that in commercial powder. This jump in the voltage-induced hole current behavior signifies the main role of the formation of interfacial defects in the nanocrystalline charged CuI clusters by mechanical grinding. Ten percent Ag-doped CuI shows a single current step at 0.29 V. The exact origin of the interesting current steps is under investigation. In the case of 15% Ag-doped CuI solid solution, a diode-like I – V curve is observed with current varying in the range of 1028 –1026 A without any current step. Twenty-five percent Ag-doped CuI also shows a similar behavior with very small slope in the range 0 # V # 0:4 but with the current one order less than the 15%
Fig. 11. I – V characteristics for both pure CuI and Cu-rich solid solutions at 300 K.
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Ag-doped CuI. Notably, we could observe the increases in the threshold voltage with the substitution of Ag into CuI. Also, at the same time the order of the magnitude of the current decreases from 1023 to 1027 A in the Cu-rich solid solutions as Ag is increased from 5 to 25%. 4. Conclusion Ag-rich (Ag12xCuxI) and Cu-rich (Cu12yAgyI) solid solutions and the end members AgI, CuI in the AgI –CuI system have been synthesized as nanoparticles by a soft MCR at ambient temperature. A monophasic solid solution with zincblende structure (a ¼ 637 pm) and crystallite size of about , 31 nm was realized for x ¼ 0:25 (Ag0.75Cu0.25I). With increasing Cu (doping element) concentration, the lattice parameter decreases linearly from 649 to 604 pm in accordance with Vegard’s law, signaling the development of static cation disorder in the AgI –CuI system. Stacking faults and grain boundaries formed during attrition assist in the formation of size-reduced (, 13 nm) nanocrystallites of Cu0.75Ag0.25I. The stronger covalent bond in CuI generates maximum strain in the middle (50 – 50) composition ðx; y ¼ 0:5Þ leading to the successful synthesis of unagglomerated uniform sized (, 17 nm) and spherical nanocrystallites of Ag0.50Cu0.50I (a ¼ 626 pm), which has potential for sensor applications. MCR of (Ag – Cu) I extends the sphalerite – bcc phase transition temperature ðTt Þ in AgI by strengthening the Ag –I bond in the Ag-rich region over which g-phase is stable. Cu-induced modification of cation disorder in AgI is evidenced by enhancement of transition temperature and of electronic conductivity with the latter exhibiting a diode type I – V characteristic possibly due to interfacial microscopic p –n junctions created by MCR process. MCR seems to be a unique technique to stabilize these and other ternary metastable superionic conductors.
Acknowledgements The authors sincerely thank Dr P. Senthil Kumar for helpful discussion.
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References [1] C. Suryanarayana, E. Ivanov, V.V. Boldyrev, Mater. Sci. Eng. A304– 306 (2001) 151 –158. [2] V.V. Boldyrev, 10th International Symposium, Reactivity of Solids, Dijon, 1984. [3] J.M. Criado, C. Real, 10th International Symposium, Reactivity of Solids, Dijon, 1984. [4] S. Chandra, Superionic Solids: Principles and Applications, North Holland, Amsterdam, 1981. [5] T. Takahashi, S. Ikeda, O. Yamamota, J. Electrochem. Soc. 119 (1972) 477. [6] H. Hayashi, J. Phys. Soc. Jpn 51 (1982) 811. [7] M. Kobayashi, S. Ono, T. Kohda, H. Iyetomi, S. Kashida, T. Tomoyose, Solid State Ionics 154–155 (2002) 209. [8] A. Yoshiasa, A. Inaba, T. Ishii, K. koto, Solid State Ionics 79 (1995) 67. [9] T. Sakuma, J. Phys. Soc. Jpn 57 (1988) 565. [10] Y. Wang, L. Huang, H. He, M. Li, Physica B 325 (2003) 357. [11] P. Senthil Kumar, C.S. Sunandana, Nano Lett. 2 (2002) 431. [12] J.R.G. Patnaik, Y.L. Saraswathi, C.S. Sunandana, Proc. Int. Conference on Trends in Mechanical Alloying, Oxford and IBH Publishing Co., New Delhi, India, 2002, pp. 12 –21. [13] J. Kimura, T. Ida, M. Mizuno, K. Endo, M. Suhara, K. Kihara, J. Mol. Struct. 61–69 (2000) 522. [14] P.S. Kumar, P. Balaya, P. Goyal, C.S. Sunandana, J. Phys. Chem. Solids 64 (2003) 961. [15] A.J.C. Wilson, X-Ray Optics, second ed., Lone, Methuen, 1962. [16] B.D. Cullity, Elements of X-Ray Diffraction, Addison-Wesley Publishing Co., Philippines, 1978. [17] N.W. Ashcroft, N.D. Mermin, Solid State Physics, Holt-Saunders International Editions, Philadelphia, 1976. [18] J. Rogez, O. Schaf, P. Kanauth, Solid State Ionics 955 (2000) 136. [19] M.W. Barsoum, Fundamentals of Ceramics, McGraw Hill, New York, 1987, Chapter 4, p. 95. [20] R.W. Vook, F. Witt, J. Vac. Sci. Technol. 2 (1965) 49. [21] P. Babu Dayal, B.R. Mehta, Y. Aparna, S.M. Shivaprasad, Appl. Phys. Lett. 81 (2002) 4254. [22] J.I. Langford, J. Appl. Crystallogr. 4 (1971) 164. [23] K. Wakamura, Solid State Ionics 149 (2002) 73. [24] M.J. Rice, S. Strassler, G.A. Tombs, Phys. Rev. Lett. 32 (1974) 596. [25] M.A. Korzhuev, Phys. Solid State 40 (2) (1998) 204 –205. [26] P. Senthil Kumar, Y.L. Saraswathi, C.S. Sunandana, Mater. Phys. Mech. 4 (2001) 71. [27] M.M. Hebb, Chem. Phys. 20 (1952) 185. [28] C. Wagner, J. Electrochem. Soc. 60 (1956) 4.
Nanophases in mechanochemically synthesized AgI– CuI system: structure, phase stability and phase transitions D. Bharathi Mohan, C.S. Sunandana* School of Physics, University of Hyderabad, Hyderabad 500 046, India Received 20 October 2003; revised 24 March 2004; accepted 9 April 2004
Abstract Nanoscale crystallites of Ag-rich (Ag12xCuxI, x ¼ 0:05; 0.10, 0.15 and 0.25), Cu-rich (Cu1-yAgyI, y ¼ 0:05; 0.10, 0.15 and 0.25) and intermediate Ag1-xCuxI ðx ¼ 0:50Þ solid solutions and end members AgI, CuI with sizes in the range of 46 –13 nm were synthesized by attrition at ambient temperature in a soft mechanochemical reaction (MCR) of Ag, Cu and I. Monophasic g-AgI (zincblende, a ¼ 638 pm) with disordered Agþ sublattice and the crystallite size of about , 31 nm was realized in the case of Ag0.75Cu0.25I ðx ¼ 0:25Þ composition. Lattice parameter decreases linearly from 649 to 604 pm with increasing Cu concentration in the AgI– CuI system validating Vegard’s law. Smallest size (,13 nm) agglomerated nanocrystals were realized in the Cu-rich composition Cu0.75Ag0.25I (a ¼ 615 pm), while unagglomerated uniform-sized (,17 nm) and spherical shape nanocrystallites of Ag0.50Cu0.50I (a ¼ 626 pm) with maximum strain were synthesized for sensor applications using MCR. Differential scanning calorimetry study shows the systematic changes in the phase transition temperature with Cu substitution. Ag-rich composition posses less enthalpy (DH (x or Cu ¼ 0.05, 0.10, 0.15, 0.25) ¼ 6.0, 6.11, 6.6, 6.3 in kJ/mol) and entropy (DS (y or Ag ¼ 0.05, 0.10, 0.15, 0.25) ¼ 14.15, 14.1, 15.03, 13.6 in J/mol K) when compared to undoped AgI (DH ¼ 9:63 kJ=mol; DS ¼ 22:8 J=mol K) implying greater thermal stability of g-phase due to Cu-strengthened Ag– I bond. Enhanced entropy (DS ¼ 8:17 J=mol K) in Cu0.75Ag0.25I (Cu-rich) solid solutions relative to CuI (DS ¼ 1:0 J=mol K) indicates Ag-induced cation disorder. Fifteen percent Ag-doped CuI (Cu0.85Ag0.15I) nanocrystals apparently behave like microscopic p – n junctions with currents in the range of 1026 –1028 A characterized by a non-linear I – V curve. q 2004 Elsevier Ltd. All rights reserved. Keywords: D. Phase transitions
1. Introduction Mechanochemical reaction (MCR) or mechanical grinding has recently emerged as a viable and convenient method to produce nanoscale ceramics [1]. MCR triggered by the application of mechanical energy, is characterized by a large negative fractional energy change and is thermodynamically feasible at ambient temperature. The over all reaction can be controlled by the mechanical generation of clean/fresh surfaces due to fracturing, increased defect density and reduction in particle size. Boldyrev [2] described the main methods of investigation of the mechanochemical process in inorganic solids, as well as the influence of preliminary mechanical processing on the reactivity of * Corresponding author. Tel.: þ 91-40-23134324; fax: þ 91-4023010227. E-mail address: [email protected] (C.S. Sunandana). 0022-3697/$ - see front matter q 2004 Elsevier Ltd. All rights reserved. doi:10.1016/j.jpcs.2004.04.007
solids. Grinding generally influences texture and structure leading to a decrease of the particle size of the product phase and, simultaneously, to an increase of the microstrain due to the contribution of the grain boundaries formed during the process [3]. AgI (CuBr and CuI) [5,6] are well-studied superionic conducting materials with zincblende structure at room temperature, and rather high Agþ conductivity [4]. The p –d hybridized [7] Ag – I [8] and Cu – I [9] bonds are weakly tetrahedral, with ease of dislocations/stacking faults production in the rather soft sphalerite or zincblende lattice, with cation disorder. This unique feature is conducive for realizing metastable and new superionic phases by mechanochemistry. CuI is a mixed ionic– electronic conductor at room temperature with predominant electron –hole conduction up to 200 8C. (1) The cubic (superionic) a-AgI forms above 146.5 8C, in which iodine ions form a bcc lattice and two silver ions are distributed randomly among the many available crystallographic sites. (2) b-hexagonal
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(non-superionic) phase (Wurtzite structure) of AgI forms reversibly upon cooling the high temperature cubic form below 146.5 8C in which each iodine ion being surrounded by a regular tetrahedron of four silver ions and vice versa. (3) g-AgI, with a disordered Agþ sublattice in its zincblende structure and possessing an Agþ conductivity higher than that of b-AgI, is metastable below 137 8C [4,10]. An ambient solid-state process such as the MCR, when applied to these ‘soft’ superionic conductors AgI and CuI, can be used to synthesize new materials in terms of both composition and microstructure applicable in ‘as prepared’ form in electrochemical sensor and battery systems at ambient. In this paper, the influence of mechanical work on the constituents of Ag-rich (Ag12xCuxI) and Cu-rich (Cu12yAgyI) solid solutions and the end members AgI, CuI has been systematically investigated in a MCR using XRD, scanning electron microscope (SEM), differential scanning calorimetry (DSC) and electronic conductivity as probes. A new phase (Ag0.50Cu0.50I) has been obtained using this unique but convenient technique. A variety of elementary physicochemical processes at micro- and macro-levels are expected to be stimulated leading to such metastable nanocrystalline products with optimal (micro) structure and morphology.
2. Experimental methods Compositions corresponding to AgI, Ag-rich (Ag12xCuxI where x ¼ 0:05; i.e. 5%, 0.10, 0.15, 0.25), intermediate Ag12xCuxI ðx ¼ 0:50Þ; Cu-rich (Cu12yAgyI, y ¼ 0:05; 0.10, 0.15, 0.25) portion of the AgI –CuI system and CuI were synthesized by mechanical grinding in a 600 agate mortar and pestle for 5 h at room temperature in an unilluminated room by using appropriate quantities (in wt%) of copper (LOBA, India), silver (Special materials project, Hyderabad, India) and I (Rasayana Laboratory, India). When attrition was carried out, the charge (Ag or Cu or I or Ag þ Cu or Cu þ Ag or Ag þ Cu þ I or Cu þ Ag þ I) present at the bottom of the mortar is sheared by the pestle at the start of the spiral type motion executed by it. This motion causes the charges to spread over the inner surface of the mortar; say 1/3 of the inner surface area (294.5 cm3). In about a minute of attrition the charge undergoes ‘reactive mixing’ as the pestle circles about 60 times in the mortar—a motion that is fairly uniform and periodic. The periodic motion of the pestle causes Ag/Cu grains to get compressed and the compressed surface is exposed to iodine molcules on the breaking-up and sublimating iodine flakes. Structural characterization on these samples was performed using a PHILIPS X-Ray diffractometer with Cu Ka1 radiation. Microstructure studies were carried out using PHILLIPS XL-30 SERIES SEM by spreading the assynthesized powder over a well-cleaned microscopic slide after applying vacuum grease. Phase stability and phase transitions were investigated using a DSC instrument (TA, USA) with crimped samples held in an argon atmosphere.
The powdered sample was compressed at pressures of 4 ton in a stainless steel dye at room temperature. The pellets thus obtained had an area of 8 mm and thickness of 3 mm. Pellets were annealed at 373 K(^ 1 K) for 3 h and used for measuring I – V characteristics. I – V measurements were made by a two-probe method. In order to measure the electronic conductivity, a graphite electrode was used to block the ion movement on the other side of the sample configuration. The electrode configuration in this case is M/MX/B where M ¼ Ag or Cu, MX ¼ (Ag, Cu)I, B ¼ graphite. A constant voltage source was connected in series with the sample and a fixed (standard) resistance of 10 ohm. A KEITHLEY 195A digital multimeter was used for measuring the current flow through the sample.
3. Results and discussion 3.1. Structure of b- and g-AgI and CuI, Vegard’s law and particle size Fig. 1 (expanded spectrum includes only major lines) shows the XRD patterns of as prepared AgI –CuI solid solutions. Undoped AgI sample (pattern (a)) contains two phases: the major g-AgI (sphalerite or zincblende structure, approx. 85.55%) phase, characterized by the XRD peaks (111), (220), (311), (400), (331), (422) and (511) [a ¼ 649 pm], and the minor b-AgI (Wurtzite structure, appox. 14.45%) identified by (100), (101), (102) and (103) reflections [a ¼ 412 pm; c ¼ 730 pm]. The main purpose of Cu doping is to stabilize the metastable zincblende structure
Fig. 1. X-ray diffractograms of AgI–CuI solid solutions mechanically ground for 300 min. (a) AgI; (b) Ag0.95Cu0.05I; (c) Ag0.90Cu0.10I; (d) Ag0.85Cu0.15I; (e) Ag0.75Cu0.25I; (f) Ag0.50Cu0.50I; (g) Cu0.75Ag0.25I; (h) Cu0.85Ag0.15I; (i) Cu0.90Ag0.10I; (j) Cu0.95Ag0.05I; and (k) CuI.
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of g-AgI besides assisting in particle size reduction. g-AgI nanoparticles were clearly seen upon progressively iodized Ag – Cu metastable alloy thin films under ambient conditions, which have been studied, by optical absorption and photoluminescence [11]. The Ag0.95Cu0.05I ðx ¼ 0:05Þ sample (pattern (b)) has two b-AgI lines (100) and (104) [approx. 3.6% in relative intensity] and one silver line (111) with six g-AgI lines (approx. 96.4%). Ag0.90Cu0.10I ðx ¼ 0:10Þ composition has only one b-AgI line (100) with negligibly small intensity (approx. 1.8%) [pattern(c)]. The Bragg reflections of Ag0.85Cu0.15I sample [pattern (d)] show predominant peaks corresponding to the zincblende structure of g-AgI alone. More significantly, the growth of wurtzite phase of b-AgI is now completely suppressed, which suggests a rapid nucleation of g-AgI crystallites aided by Cu. Moreover, it contains unreacted Ag (111), signaling the formation of monophasic g-AgI. But in the case of Ag0.75Cu0.25I [pattern (e)], there is no Ag line implying complete consumption of Ag, so that the pattern corresponds to monophasic g-AgI with sharper peaks. Fig. 2 shows a gradual increase of all the Bragg angles accompanied by a broadening of the (111) plane conspicuously observed upon progressive Cu addition. The lattice parameter (a in pm) decreases linearly with increasing Cu (wt%) in a continuous manner (Fig. 3). As more and more Cu atoms substitute Ag in AgI, the Ag – I bond progressively shortens, there by stabilizing the g-AgI structure. In the Cu-rich region, Ag is accumulated in the already stable zincblende CuI lattice. The systematic development and stabilization of g-CuI using grinding time as a parameter have already been discussed [12]. Undoped CuI [pattern (k)] shows clearly, the more fully developed Bragg reflections of g-CuI lines
Fig. 2. Gradual increase of Bragg angle accompanied by a broadening of the (111) plane upon progressive Cu addition. (a) AgI; (b) Ag0.95Cu0.05I; (c) Ag0.90Cu0.10I; (d) Ag0.85Cu0.15I; (e) Ag0.75Cu0.25I; (f) Ag0.50Cu0.50I; (g) Cu0.75Ag0.25I; (h) Cu0.85Ag0.15I; (i) Cu0.90Ag0.10I; (j) Cu0.95Ag0.05I; (k) CuI; and (l) commercial powder of CuI.
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Fig. 3. Plot of lattice parameter ðaÞ vs Cu composition in the AgI– CuI system of solid solutions showing smooth linear decrease of ‘a’ from AgI to CuI.
characterized by (111), (200), (220), (311), (222), (400), (331), (420) and (422) with intensities larger than those of Cu-rich (Cu12yAgyI) phases. A continuous shift in the Bragg angle is observed in all Bragg peaks [patterns (j), (i), (h) and (g)] of Cu-rich system as it moves towards Ag-rich region (Fig. 1) upon progressive addition of Ag. From Fig. 3, we observe a monotonic linear increase in the lattice parameter (a in pm) of all compositions relative to undoped CuI. The smaller ionic size of Cuþ (145 pm) as compared to Agþ (165 pm) probably results in the systematic increase in the lattice parameter (a) in a linear fashion with increasing Ag addition [13,14]. The Bragg reflections of the intermediate (50:50) composition Ag0.50Cu0.50I (where x ¼ 0:50) lie in-between Ag0.75Cu0.25I (pattern (e)) and Cu0.75Ag0.25I (pattern (g)). The changes in 2u of the first three prominent peaks are, 24.38 (111); 40.41 (220); 47.77 (311). Also the intensities of (200), (222) and (420) lines become negligibly small. The smooth decrease in ‘a’ (pm) (or unit cell volume) of AgI is linear as Cu concentration is increased (Fig. 3). This indicates a continuously uniform stabilization of static cation disorder in the sphalerite or zincblende AgI lattice brought about by the mechanochemical processing. Finally, Vegard’s law seems to be valid in AgI – CuI system [15]. All samples produce appreciable diffraction broadening and it is reasonably assumed that this arises from small crystallites and internal stresses. Diffraction theory predicts that the breadth due to crystallite size varies with angle as sec u and that due to elastic strain as tan u [15]. The additional broadening in diffraction peaks beyond the inherent peak widths due to instrumental effects can be used to measure crystallite sizes as low as 1.0 nm. The crystallite sizes of as-prepared compositions were calculated from the full width at half maximum (FWHM) of three prominent peaks (111), (220) and (311) using the Debye –Scherrer formula, t ¼ ½ð0:9lÞ=ðB cos uÞ [16].
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Use of the Lorentzian function to fit the FWHM of three peaks yields more accurate (error is approx. 0.22%) results than those obtained by using Gaussian (1%) and Voigt functions (0.5%). Therefore, the observed XRD lines are essentially Lorentzian, pointing to a uniform crystallite size distribution. These apparent sizes essentially correspond to the true size of the crystallites, because the observed induced microstrain is generally very small (, 1023 – 1024) due to the relatively small mechanical energy input to the system during the MCR process. Furthermore, the elastic constants (particularly the shear elastic constant (C44)) of Ag (0.46 £ 1012 dynes/cm2) and Cu (0.75 £ 1012 dynes/ cm2) at 300 K [17] are very close and are relatively small, low power, soft attrition process does not induce much strain in the fully reacted solid solution samples. Overall, we can say that the broadening is due to the microstrain presence in the crystallites is negligible. Effects of Cu addition in reducing the particle size in the Ag12xCuxI solid solutions prepared by aqueous solution method in our laboratory were discussed previously [14]. Fig. 2 shows the gradual increase of Bragg angle accompanied by a broadening of the (111) plane upon progressive Cu addition in the Ag-rich region. Fig. 4 clearly shows the crystallite size vs Cu composition plot for the AgI –CuI system. Bigger size (, 46 nm) crystallites were observed in undoped AgI, because it contains both b-AgI (, 14.5%) and g-AgI (, 85.5%) phases. 5% Cu addition shows a sudden fall in the crystallite size (, 31 nm). The smaller Cuþ ions are expected to be present mostly in the subsurface region of the crystals rather than in the crystal bulk, facilitating further reduction in the interfacial energies between the randomly distributed polytypes of AgI, leading to the stabilization of the metastable cubic zincblende g-AgI with smaller crystallites [18]. Interestingly, the smallest crystallite size namely,
Fig. 4. Crystallite size changes in AgI–CuI solid solutions made by mechanochemical reaction. Smallest crystallite size occurs for Cu0.75Ag0.25I but unagglomerated crystallites of 17 nm are seen in Ag0.50Cu0.50I.
26 nm is observed in the Ag-rich region, which corresponds to the Ag0.90Cu0.10I solid solution. It is significant that Ag0.75Cu0.25I ðx ¼ 0:25Þ solid solution contains slightly bigger crystallites of about 31 nm, due to the absence of strain that is normally induced by the MCR process. It is interesting to look at the strain analysis plot where x ¼ 0:25 of Cu-doped system shows a horizontal straight line (see later) in contrast to undoped AgI and Ag-rich (x ¼ 0:05 and 0.10 Cu doped) compositions. In the case of Cu-rich samples, increase in Ag addition leads to a systematic broadening of the XRD peaks (Fig. 2 shows the (111) plane). So the particle size also starts decreasing from undoped CuI (, 31 nm). The Cu-rich composition Cu0.75Ag0.25I ðy ¼ 0:25Þ shows the broadest peak, corresponding to the smallest particle size (, 13 nm) observed, due to the presence of strain, stacking faults and grain boundaries. This implies that shear strain and defects together reduce the size of the nanoparticles synthesized by the MCR process. It is interesting that even 5% Cu substituted into AgI brings about a considerable reduction in the particle size. This can explained by noting that the polarizing power of the smaller-sized Cuþ is greater than that of Agþ [19]. Also the Cu – I bond is more covalent than the Ag – I bond. For these two reasons, the presence of even a small amount ($ 5%) of Cu in AgI initiates the particle size reduction processes, one of which may be the considerable reduction of forces responsible for agglomeration. So, regardless of whether one is dealing with Cu-rich or Ag-rich compositions, the particle size reduction invariably seems to take place. 3.2. Strain analysis The Nelson – Reily function (NRF) vs Dd=d curves [20,21] and least squares regression line: ð2wÞ2 cos2 u as a function of sin2 u [22] for AgI –CuI system obtained from the XRD data are plotted in Fig. 5a,b, respectively. Systematic lattice transformation of all crystal planes upon Cu substitution is clearly evident (Fig. 5a). Dd=d values increase up to x ¼ 0:50 and thereafter it starts decreasing all the way down to undoped CuI ðy ¼ 0Þ: In the case of Ag-rich, Cu-rich (y ¼ 0:05; 0.10) and undoped CuI samples, Dd=d values lie on a horizontal straight line, almost parallel to NRF axis, which indicates the absence of stacking faults in the cubic structure. But, rather large scatter observed in Dd=d values for the intermediate 50– 50 (Ag0.50Cu0.50I) and Cu-rich (Cu0.75Ag0.25I, Cu0.85Ag0.15I) samples is a measure of the high concentration of stacking faults present in these compositions. Ag-rich and Cu-rich samples show less scatter with a small slope value whereas intermediate (50, 75% Cu doped) compositions gives large slope value (Fig. 5b). Average microstrain is calculated and tabulated for the all compositions by measuring slopes of the least squares regression line (Table 1). Maximum strain values are observed only in the intermediate 50 and 75% Cu
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Fig. 5. (a) Strain analysis of AgI–CuI solid solutions obtained from XRD. Strain analysis of AgI–CuI solid solutions obtained from XRD. (A) AgI; (B) Ag0.90Cu0.10I; (C) Ag0.75Cu0.25I; (D) Ag0.50Cu0.50I; (E) Cu0.75Ag0.25I; (F) Cu0.85Ag0.15I; (G) Cu0.95Ag0.05I; see text for details. (b) Least square fit analysis shows the induced microstrain obtained from XRD. Slope gives the average strain value. (a) Ag0.75Cu0.25I; (b) Ag0.50Cu0.50I; (c) Ag0.10Cu0.90I; and (d) Ag0.85Cu0.15I. Other compositions (not marked) do not show significant strain.
doped AgI compositions but not in others. The larger covalency of CuI (relative to AgI) [23] possibly assists in generating considerable strain in Cu-rich compositions leading to size-reduced and unagglomerated Ag0.50Cu0.50I crystallites as seen in SEM. This is in conformity with the results of particle size analysis where, these compositions have minimum particle size. Less scattered points are observed in pure AgI system due to the presence of hexagonal (b-AgI) structure. There are two stages of structural transformation observed in the presence of stacking faults in this system one being the transformation of the mixture of minor Wurtzite (b-AgI) and major zincblende (g-AgI) phase to pure zincblende (g-AgI) phase and the other being the pure g-AgI phase to zincblende (g-CuI) phase transition. The latter one is essentially a weakly covalent to a strongly covalent bonding transition taking place gradually in the Cu-rich phases. 3.3. Uniform size MCR Ag0.50Cu0.50I crystallites (13 nm) Secondary electron images show qualitatively the differences in the surface morphology of mechanochemically reacted ground samples of AgI, Ag0.75Cu0.25I, Ag0.50Cu0.50I, Cu0.75Ag0.25I, Cu 0.90 Ag 0.10 I and CuI (Fig. 6). Random orientation of bigger size (, 46 nm) crystallites were observed in 300 min undoped AgI sample (Fig. 6a). As the concentration of doping element (25%Cu) is increased, the crystallite size also decreases to , 31 nm, so that for middle composition (Ag0.50Cu0.50I), crystallites with nearly uniform size (, 17 nm) (Fig. 6c) in a completely formed solid solution were observed indicating maximized
contact areas between Ag and Cu leading to optimal reactivity. The formation of this new composition Ag0.50Cu0.50I is of both fundamental and technical interest because of the difficulty of using conventional techniques such as fusion and solid-state reaction for its realization. Agglomerated crystallites with small size (, 13 nm) and increased reactive surface area were observed in the (Cu-rich) Cu0.75Ag0.25I composition (d). Unreacted Ag (bigger particles in Fig. 6e) with agglomerated Cu particles was observed in Cu0.90Ag0.10I ðy ¼ 0:10Þ sample. Eventhough, the MCR technique is a non-equilibriumprocessing route, our SEM (and XRD) observation suggests that apart from incorporating residual strain, the samples obtained are by and large homogeneous. The relatively low energy cost for, and the high efficiency of, the process could lead to homogeneous samples. Table 1 Average apparent crystallite size and strain Compositions
Crystallite size (nm)
Microstrain (1024)
AgI Ag0.95Cu0.05I Ag0.90Cu0.10I Ag0.85Cu0.15I Ag0.75Cu0.25I Ag0.50Cu0.50I Ag0.25Cu0.75I Ag0.15Cu0.85I Ag0.10Cu0.90I Ag0.05Cu0.95I CuI
45.7 31.0 25.6 27.0 31.4 17.0 13.0 17.1 24.0 26.7 30.6
– – – 3.05 – 3.15 18.35 – 3.76 – –
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ðDSÞ—at the phase transitions are given by DH ¼ kA=m where k is the calibration constant, A is the area under the curve and m is the mass of the sample DS ¼ DH=Tt
Fig. 6. Secondary electron micrographs for MCR derived compositions in the AgI–CuI system. (a) AgI; (b) Ag0.75Cu0.25I; (c) Ag0.50Cu0.50I; (d) Cu0.75Ag0.25I; (e) Cu0.90Ag0.10I; (f) CuI. Note: uniform sized, unagglomerated nanocrystals in (c) and smallest, agglomerated nanocrystals in (d).
3.4. Thermal analysis, enthalpy and entropy changes Fig. 7 shows the thermal response of MCR derived AgI – CuI solid solutions scanned at a rate of 10 K per minute in the non-isothermal mode. The changes in thermodynamic parameters—enthalpy ðDHÞ and entropy
Fig. 7. AgI–CuI system shows the systematic variation of phase transition temperature ðTt Þ curve, where (A) AgI; (B) Ag 0.95 Cu 0.05 I; (C) Ag0.90Cu0.10I; (D) Ag0.85Cu0.15I; (E) Ag0.75Cu0.25I; (F) Ag0.50Cu0.50I; (G) Cu0.75Ag0.25I; (H) Cu0.85Ag0.15I; (I) Cu0.90Ag0.10I; (J) Cu0.95Ag0.05I; and (K) CuI.
where Tt is the transition temperature. On heating undoped AgI, a sharp and intense endothermic peak, g ! a phase transition ðTt Þ observed at 422.6 K (420 K is in commercial AgI). The increased Tt probably arises from the relatively strengthened Ag – I bond, an effect essentially induced by the MCR process. In Ag-rich region, Ag –I bond is further strengthened by the Cu substitution. So that MCR of (Ag– Cu) I extends the transition temperature ðTt Þ range over which the g-phase is stable (Fig. 8). This extended thermal stability of the g-phase is important in view of the potential applications. A competition between Cu-induced surface area enhancement and thermally induced surface area reduction may lead to an increased cation-sublattice melting temperature reflected in DSC curves. AgI is cation disordered even at room temperature and also a weakly covalent bonded structure. The addition of Cuþ has perhaps resulted in a considerable decrease in the number of Frenkel defects with the maximum strain field abruptly produced at the g – a phase transition leading to a broad and increased asymmetric endothermic peak that systematically develops in x ¼ 0:10; 0.15, 0.25 (Ag-rich) solid solutions. It is the interaction of the interstitial cation disorder with the strain field associated with this net volume change which is the principal factor for the Cu concentration-dependence of the phase transition temperature ðTt Þ [24]. Thus, grinding causes local changes in temperature and free energy that are tied up with the crystallite size reduction. ‘Melting’ of the sublattice of mobile ions generally occurs at the temperature of the superionic
Fig. 8. AgI–CuI system shows the changes in the phase transition temperature ðTt Þ. Whereas pure AgI (at 0 wt%), Ag-rich and Cu-rich compositions exhibit g ! a transition and pure CuI (at 100 wt%) exhibits both g ! b and b ! a transitions.
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phase transition when some of the SIC ions are suddenly found in intersitial sites (Frenkel defects), where they are distributed over many vacancies (order –disorder phase transition). This sudden increase in the Frenkel defect concentration contributes significantly to the total entropy of the superionic transition. The entropy of the crystallization of materials such as Cu, Ag, Au, and Li in a ‘molten’ sublattice is explained by Korzhuev [25]. Fig. 9 shows the enthalpy/entropy vs Cu substitution in weight % at the phase transition temperature ðTt Þ: Undoped AgI shows the large value of enthalpy and entropy (DH ¼ 9:63 kJ=mol; DS ¼ 22:8 J=mol K) due to the presence of large number of Frenkel defects induced by the MCR process. Even a 5% addition of Cu causes a sudden fall (followed by a more gradual decrease) in the values of enthalpy (DH (x or Cu ¼ 0.05, 0.10, 0.15, 0.25) ¼ 6.0, 6.11, 6.6, 6.3 in kJ/mol) and entropy (DS (x or Ag ¼ 0.05, 0.10, 0.15, 0.25) ¼ 14.15, 14.1, 15.03, 13.6 in J/mol K) observed. These observations imply that the Cuþ ion ‘arrests’ the order –disorder phase transition in AgI, which means that, the number of Frenkel defects is decreased in the Ag-rich solid solutions. For the middle Ag0.50Cu0.50I composition, a single-phase transition is observed at 548.6 K, which is to be compared to 420 and 648 K (g ! b transition). The Cu –I bond is more covalent than the Ag –I bond and the strength of the former is 55 meV compared to 36 meV for the latter. On heating 5-h ground undoped CuI, two endothermic peaks are observed, one at 648 K corresponding to g ! b phase transition and the other at 671.7 K representing the b ! a phase transition of CuI. Increased Tt (DT < 6 K) in the g ! b phase transition and decreased Tt (DT < 8 K) in the b ! a phase transition are observed when compared to the phase transitions in the commercial powder CuI. This suggests that while the MCR process strengthens the g-CuI (zincblende, F4¯3m) structure at the same time it weakens the a-CuI (zincblende, Fm3¯m)
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structure. Attrition essentially prepares the fcc Cu lattice so that it favors the accommodation of I ions in a weakly covalent arrangement [26]. The undoped CuI shows very less enthalpy (DH ¼ 0:5 kJ=mol) and entropy (DS ¼ 1:0 J=mol K) compared to Cu-rich iodide solid solutions (see below). As 5% Ag is added to CuI, only one peak is observed at 649.7 K, which corresponds to g ! a phase transition implying that the g ! b transition is completely suppressed. The occurrence of a single-phase transition in CuI probably means that the difference between the structures of the b-phase and the g-phase has had no significant effect on the nature of the phase transition. As the Cu concentration is increased, a continuous decrease in the phase transition temperature was observed (Fig. 8). Also, there is no trace of the g ! a characteristic of AgI phase transition. Therefore, all Agþ ions have apparently been incorporated into the CuI lattice, partially replacing Cuþ ions in the cation-sublattice. The effects of Ag in Cu-rich solid solutions are to cause an increase in the surface energy of the nanoparticle and to induce an increase in the cation disorder in the Cuþ-ion sublattice. Consequently, larger values of enthalpy (DH (y or Ag ¼ 0.05, 0.15, 0.25) ¼ 5.8, 6.23, 4.917 in kJ/mol) and entropy (DS (y or Ag ¼ 0.05, 0.15, 0.25) ¼ 8.9, 9.95, 8.17 in J/mol K) are observed in Cu-rich solid solutions (Fig. 9). Phase transitions are not multiply reversible implying ‘thermal annealing’ of strains induced by the MCR process. The kinetics of phase transitions and the structural relaxation of MCR samples of AgI –CuI solid solutions are interesting problems that deserve to be studied separately in depth. 3.5. I – V curve The process of MCR introduces a number of steps in the formation of solid solutions; the important one being charged interface formation. These interfaces probably act as microscopic p – n junctions, and the process eventually produce a non-negligible concentration of electronic current carriers in the MCR derived nanocrystalline powder. Cu-rich compositions tend to contain holes and Ag-rich compositions contain electrons as current carriers in the I–VII mixed semiconductors (while AgI is an n-type semiconductor, CuI is a p-type semiconductor). To check if MCR samples indeed possess non-linear I – V characteristics, a dc two-probe Hebb– Wagner [27,28] polarization cell of the type, ð2Þ AglAg12x Cux IlC ðþÞ; ð2Þ CulCu12x Agx IlC ðþÞ;
Fig. 9. AgI–CuI system shows the variation of Enthalpy (V) and entropy (†) in both heating (endotherm) and cooling (exotherm) cycles.
was constructed. When a dc potential (V) less than the decomposition potential of the electrolyte is applied with positive polarity on the right side, the Agþ ions move initially from the right to Ag electrode. Graphite acts as a blocking electrode so that there is no supply of Agþ ions at the ClAgrich interface. At the steady-state, the potential gradient due to electric field is just balanced by that due to the chemical
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potential gradient so that the ionic current is zero. Under these conditions, the current is solely due to differential flow of electrons or holes. Wagner has shown that the steady-state current (I) is given by I ¼ Ie þ Ih ¼ ðkTA=eLÞ½se {1 2 expð2eV=kTÞ} þ sh {expðeV=kTÞ 2 1} where se ; sh are the conductivity due to electrons and holes, respectively. L is the length, A the area of the sample, k; the Boltzman constant and V is the potential. Ag-rich solid solutions show the systematic variation in the I – V curve (Fig. 10a). Pure AgI shows the linear curve with the trivial current flow in the order of 1029 –1028 A. There are no significant effects observed up to 10% Cu substitution. The Ag0.85Cu0.15I solid solution shows a ramp-like current behavior up to 0.35 V and nonlinearity just shows up thereafter. The sudden rise in the current observed after 0.35 V in the order of 1028 –1027 A in a non-linear manner in the Ag0.75Cu0.25I solid solution,
Fig. 10. (a) I – V characteristics of pure AgI and Ag-rich solid solutions at 300 K. (b) I – V characteristics of intermediate 50–50 solid solution at 300 K.
suggests that a substantial number of holes are injected by Cu in the Ag-rich solid solutions. The current behavior was found to be different for the two electrodes Ag and Cu used in the experiments. Fig. 10b shows the non-linear response of current with increasing voltage for the Ag0.50Cu0.50I composition. Though, there is no noticeable change in the current behavior up to V < 0:30 V; there is a significant difference between the characteristic V – I behavior of the sample with Ag as electrode and with Cu as electrode after V $ 0:35 V: While, it is the residual current behavior below V < 0:30 V; the change in current behavior at the threshold voltage of 0.35 V implies that the voltage-induced hole current response similar to that of a forward biased p –n junction. Fig. 11 shows the electronic conductivity for pure CuI and Cu-rich solid solutions. On the other hand, CuI (commercial powder) gives a nearly linear curve with the current changing from 1026 to 1024 A. But there are two current steps found in the MCR sample CuI: one at < 0.2 V and the other at 0.33 V. Thereafter, it is less linear than I – V plot of the commercial powder. The obtained current magnitude is one order higher than that in commercial powder. This jump in the voltage-induced hole current behavior signifies the main role of the formation of interfacial defects in the nanocrystalline charged CuI clusters by mechanical grinding. Ten percent Ag-doped CuI shows a single current step at 0.29 V. The exact origin of the interesting current steps is under investigation. In the case of 15% Ag-doped CuI solid solution, a diode-like I – V curve is observed with current varying in the range of 1028 –1026 A without any current step. Twenty-five percent Ag-doped CuI also shows a similar behavior with very small slope in the range 0 # V # 0:4 but with the current one order less than the 15%
Fig. 11. I – V characteristics for both pure CuI and Cu-rich solid solutions at 300 K.
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Ag-doped CuI. Notably, we could observe the increases in the threshold voltage with the substitution of Ag into CuI. Also, at the same time the order of the magnitude of the current decreases from 1023 to 1027 A in the Cu-rich solid solutions as Ag is increased from 5 to 25%. 4. Conclusion Ag-rich (Ag12xCuxI) and Cu-rich (Cu12yAgyI) solid solutions and the end members AgI, CuI in the AgI –CuI system have been synthesized as nanoparticles by a soft MCR at ambient temperature. A monophasic solid solution with zincblende structure (a ¼ 637 pm) and crystallite size of about , 31 nm was realized for x ¼ 0:25 (Ag0.75Cu0.25I). With increasing Cu (doping element) concentration, the lattice parameter decreases linearly from 649 to 604 pm in accordance with Vegard’s law, signaling the development of static cation disorder in the AgI –CuI system. Stacking faults and grain boundaries formed during attrition assist in the formation of size-reduced (, 13 nm) nanocrystallites of Cu0.75Ag0.25I. The stronger covalent bond in CuI generates maximum strain in the middle (50 – 50) composition ðx; y ¼ 0:5Þ leading to the successful synthesis of unagglomerated uniform sized (, 17 nm) and spherical nanocrystallites of Ag0.50Cu0.50I (a ¼ 626 pm), which has potential for sensor applications. MCR of (Ag – Cu) I extends the sphalerite – bcc phase transition temperature ðTt Þ in AgI by strengthening the Ag –I bond in the Ag-rich region over which g-phase is stable. Cu-induced modification of cation disorder in AgI is evidenced by enhancement of transition temperature and of electronic conductivity with the latter exhibiting a diode type I – V characteristic possibly due to interfacial microscopic p –n junctions created by MCR process. MCR seems to be a unique technique to stabilize these and other ternary metastable superionic conductors.
Acknowledgements The authors sincerely thank Dr P. Senthil Kumar for helpful discussion.
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