Effect of Local Disorder on the Transport Properties of Al-Doped SmBa

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Effect of Local Disorder on the Transport Properties of Al-Doped SmBa2Cu3O6+δ Superconductors Marco Scavini,*,† Mauro Coduri,† Mattia Allieta,† Luca Mollica,†,‡ Michela Brunelli,§ Lorenzo Malavasi,| Alessandro Lascialfari,⊥,#,∇ and Claudio FerreroO Dipartimento di Chimica Fisica ed Elettrochimica, UniVersita` di Milano, Via C. Golgi 19, I-20133 Milano, Italy, Institut Laue-LangeVin, 6, rue J. Horowitz, BP 156, 38042 Grenoble Cedex 9, France, IENI-CNR, INSTM and Dipartimento di Chimica Fisica “M. Rolla”, UniVersita` di PaVia, Viale Taramelli 16, I-27100 PaVia, Italy, Dipartimento di Scienze Molecolari Applicate ai Biosistemi, UniVersita` degli studi di Milano, I-20122 Milano, Italy, Dipartimento di Fisica “A.Volta”, UniVersita` degli studi di PaVia, I-27100 PaVia, Italy, CNRsIstituto di Nanoscienze, I-41100 Modena, Italy, and European Synchrotron Radiation Facility, 6, rue J. Horowitz, BP 220, 38043 Grenoble Cedex, France ReceiVed: July 21, 2010; ReVised Manuscript ReceiVed: September 30, 2010

In this work we present the investigation of the structural disorder induced by Al doping in SmBa2Cu2.67Al0.33O6+δ superconductors in direct correlation with their transport properties. The interest in such superconducting materials lies in the possibility of extending the δ range of the underdoped region of REBa2Cu3O6+δ superconductors, where phenomena such as pseudogap and superconducting diamagnetic fluctuations above TC appear. The pair distribution function analysis of high-resolution powder diffraction data has shown that local distortions are present in the proximity of the Al dopant ions. Disorder spreads to the superconducting Cu2-O3 planes producing corrugations in these planes, associated with a decrease of superconducting charge carriers. Conductivity and magnetization measurements have demonstrated that Aldoping affects the charge transport mechanism: Anderson localization and superconductivity coexist, and the T and δ superconducting domains are strongly reduced. Al doping allows not only controlling chemically the charge carrier concentration but also the potential fluctuations within the Cu2-O2/3 planes and appears to be a powerful tool for the investigation of the underdoped zone of REBa2Cu3O6+δ superconductors. Introduction In recent years great investigation efforts in the field of superconductivity have been made to explore the so-called “underdoped region”, which exhibits several intriguing phenomena such as pseudogap,1,2 charge stripes,3 and superconducting fluctuations.4-8 In the case of YBCO-like superconductors (i.e., superconductors of the REBa2Cu3O6+δ family, where RE ) Y, 4f elements) underdoping can be obtained throughout suitable annealing conditions, which, in turn, determine the excess oxygen δ and the total hole density within the materials. However, the investigation of the underdoped region in this superconductor family is difficult, owing to the structural changes accompanying the increase of δ. In fact, oxidation triggers a progressive short-range ordering of Cu1-O4 ions along the crystallographic b direction, which in turn promotes a variety of long-range O orderings and phase transitions,9,10 as well as the simultaneous change of interatomic distances, due to cell parameter changes and ionic displacements. * To whom correspondence should be addressed, [email protected]. † Universita` di Milano. ‡ Present Address: Institut de Biologie Structurale, 41, rue J. Horowitz, Grenoble 38027, France. § Institut Laue-Langevin. | IENI-CNR, INSTM and Dipartimento di Chimica Fisica “M. Rolla”, Universita` di Pavia. ⊥ Universita` degli studi di Milano. # Universita` degli studi di Pavia. ∇ CNRsIstituto di Nanoscienze. O European Synchrotron Radiation Facility.

As an example, diffuse scattering studies have probed the nature and the spatial (nano-) distributions of oxygen superstructures in pure11-13 and Ca-doped13 YBCO as a function of δ. Unlike the fully oxidized YBa2Cu3O7 compound, where (ideally) infinite Cu1-O4 chains ordering is obtained along the b axis, it is quite difficult to control the chain length for the intermediate O concentrations, since, for each δ value, the O distribution within the xy (z ) 0) plane is a function of the annealing conditions (temperature T and partial pressure of oxygen pO2) and can also vary as a consequence of aging.10,14 The reduction of the total hole concentration (i.e., underdoping) can be obtained in fully oxidized samples (δ ∼ 1) also by doping the Cu site with an aliovalent cation. For this purpose Al has been chosen as a dopant to produce Al-doped SmBCO (SmBa2Cu3-xAlxO6+δ) samples. There are several reasons for this choice. First, Al3+ has a well-defined oxidation state; when an Al ion is introduced into the structure, the total hole concentration decreases by two units (taking SmBa2Cu3O6 as a reference structure15). Moreover, Al3+ is a nonmagnetic ion and substitutes for Cu only on one of the two crystallographic nonequivalent positions Cu1 (0,0,0) occupying one of the four nonequivalent (x, x, 0) x ∼ (0.06 sites.16 CuCu1 and AlCu1•• ions have different chemical environments: distorted square planar (for δ ) 1) the former and tetrahedral (without dependence on δ) the latter. Each Al ion is bound to an O4 ion in a (0, 1/2, 0) site and another one in a (1/2, 0, 0) site. When Al substitutes a Cu1 ion, its “structural role” in oxidized samples is to cut the Cu1-O4 chains into fragments the mean length of which (at fixed δ values, e.g. for δ∼1) is a function of Al concentration

10.1021/jp106805z  2010 American Chemical Society Published on Web 10/21/2010

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Figure 1. Unit cell of SmBa2Cu3O6 (left), SmBa2Cu3O7 (center), and SmBa2Cu3-xAlxO6+δ (right.) The space group for each compound is also shown. Dark-gray circles indicate full occupied oxygen sites; pale-gray circles indicate partially occupied oxygen sites. Empty squares represent empty oxygen sites for the ideal structures.

and distribution (see, e.g., Figure 10 of ref 17). For high Al concentrations the long-range structural transition (from P4/ mmm to Pmmm space group) is inhibited:18 SmBa2Cu3-xAlxO6+δ remains tetragonal even for δ ∼ 1. In Figure 1 are shown the unit cells of SmBa2Cu3O6 (left), SmBa2Cu3O7 (middle) and SmBa2Cu3-xAlxO6+δ (right). No long-range O ordering along the crystallographic b axis is allowed in the Al-doped sample. In recent papers we have showed that it is possible to control the mean Cu-O chain length of fully oxidized SmBa2Cu2.85Al0.15O∼7 samples throughout suitable annealing routes.19,20 At fixed Al and O concentrations (i.e., total hole concentration), the Cu-O ordering on the nanometric scale deeply affects the superconducting properties of the material, i.e., TC and the Meissner fraction.19 Moreover, Al clustering induces the rising of superconducting diamagnetic fluctuations above TC exhibiting unconventional Ginzburg-Landau character.8 All these experimental findings support the idea that an accurate determination of structural inhomogeneities on the nanoscale is fundamental for understanding the physical properties of strongly correlated oxides, like high TC superconductors.21 For such a reason this work aims to establish the relationship between the structural modifications induced by Al doping (both aVerage and local) in SmBa2Cu3O6+δ and the transport properties of the superconducting compound SmBa2Cu3-xAlxO6+δ, with x ) 0.33. We chose this high doping value in order to maximize the disorder induced by Al doping and to highlight its effect on conductivity and superconducting transition. The aVerage structure determined by Rietveld refinement using high-resolution X-ray powder diffraction patterns will be analyzed in the framework of the bond valence model (BVM).22 The transport properties of the Al-doped samples will be compared to those of Al-free specimens, obtained by Flor’s group.23,24 Finally, the local structural disorder introduced by doping will be invoked in order to correctly interpret the above modifications. Experimental Details Polycrystalline samples of SmBa2Cu3-xAlxO6+δ were synthesized by a solid-state reaction starting from BaO2 (Aldrich 95%), Sm2O3, CuO, and Al2O3 (all Aldrich 99.99%). Powders were mixed, pressed into pellets, and allowed to react in pure oxygen at 960 °C for a total time of 96 h with one intermediate cooling, grinding, and repelletization step. The single-phase nature of the products was controlled by X-ray powder diffraction (XRPD).

In order to produce well-defined O nonstoichiometry, aliquots of the samples were annealed under different T, pO2 conditions for 96 h or more in quartz ampules and then quenched to room T. The environmental atmosphere for the annealing step was set by flowing certified mixtures of O2 in N2, at a total pressure p ) 1 atm. Great care was taken in controlling the quenching conditions. It should be noted that all the investigated samples with the same x composition came from the same batch. Moreover, for each δ value, conductivity, susceptivity, and XRPD experiments were performed on aliquots coming from the same annealing cycle. The electrical conductivity was measured using a four-probe dc technique at both high and low T. High-temperature measurements were performed in the range 673-1073 K on a sample with parallelepiped-like shape using a Solartron 1286 electrochemical interface. The sample was placed in a horizontal furnace under gas flow at a total pressure of 1 atm. pO2 ) 1, 10-1, 10-2, and 10-3 atm were obtained using certified O2 and N2/O2 mixtures. After each measurement, T was changed steeply (typically ∆T ) 25 K) and the conductivity was measured as a function of time, in order to ensure that equilibrium with external atmosphere was reached. Low T measurements were performed on several annealed samples using the same electrochemical interface and an LTC 60 T-controller driving a Leybold ROK 19 cryostat. Using IEEE 488 interfaces, both devices were connected to a personal computer for sampling resistivity (F) versus T curves. Magnetization measurements were performed in the 2 < T < 300 K range, applying a constant magnetic field H ≈ 15 Oe, in both zero-field and field cooling conditions, by means of a standard sample extraction technique using a Quantum Design MPMS-XL7 SQUID dc magnetometer. As the field was very low, for the magnetic susceptibility χ it is possible to use the approximate formula χ ≈ M/H, where M is the measured magnetization. XRPD patterns from selected samples of SmBa2Cu3-xAlxO6+δ (x ) 0 and δ ) 0.13, 0.98; x ) 0.33 and δ ) 0.47, 0.89, 0.98) were collected at the high-resolution powder diffractometer at the ID31 beamline of the European Synchrotron Radiation Facility (ESRF), Grenoble, France.25 Samples were loaded in 0.6 mm diameter glass capillaries mounted on the diffractometer axis and spun during measurements in order to improve powder randomization; the samples were cooled down to 80 K using a

Structural Disorder Induced by Al Doping

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TABLE 1: Rietveld Refinement Results Referring to SmBa2Cu3-xAlxO6+δ Samples at 80 K (Qmax ) 18.7 Å-1)a sample space group purity (%) a/Å b/Å c/Å Uiso(Sm) z(Ba) Uiso(Ba) Uiso(Cu1/Al) z(Cu2) Uiso(Cu2) z(O1) Uiso(O1) z(O2) Uiso(O2) z(O3) Uiso(O3) Uiso(O4) f(O4) f(O5) f(Al) R(F2) RP

x ) 0, δ ) 0.13 P4/mmm 1.00 3.87500(2) ≡a 11.79863(8) 0.00082(11) 0.19275(5) 0.0035(1) 0.0063(3) 0.3561(1) 0.0013(1) 0.1514(7) 0.0070(8) 0.3739(3) 0.0045(6) ≡U(O3) 0.13 0.050 0.095

x ) 0, δ ) 0.98 Pmmm 1.00 3.83835(2) 3.89677(3) 11.6855(1) 0.00002(13) 0.18379(8) 0.00248(14) 0.0036(3) 0.3512(1) 0.00091(17) 0.1593(9) 0.0036(11) 0.3735(8) 0.0052(10) 0.3726(8) 0.0047(9) 0.004(2) 0.80(2) 0.18(2) 0.042 0.096

x ) 0.33, δ ) 0.47 P4/mmm 0.97915(5) 3.89074(1) ≡a 11.73068(6) 0.00234(9) 0.19151(5) 0.00589(11) 0.0056(3) 0.35595(8) 0.00220(10) 0.1524(6) 0.0149(11)

x ) 0.33, δ ) 0.98 P4/mmm 0.9706(7) 3.88975(1) ≡a 11.6719(1) 0.0021(1) 0.18635(7) 0.0061(1) 0.0053(3) 0.3536(1) 0.0026(1) 0.1524(8) 0.018(2)

x ) 0.33, δ ) 0.89 P4/mmm 0.9710(6) 3.89096(1) ≡a 11.6805(1) 0.0020(1) 0.18745(7) 0.0056(1) 0.0056(3) 0.3540(1) 0.0025(1) 0.1520(9) 0.022(2)

0.3739(3) 0.0060(6) 0.030(8) 0.47

0.3724(4) 0.0072(8) 0.016(3) 0.98

0.3729(4) 0.0068(8) 0.011(3) 0.89

0.33 0.063 0.080

0.33 0.068 0.080

0.33 0.071 0.076

a The following sites were considered in the structural model: Sm1(1/2, 1/2, 1/2), Ba(1/2, 1/2, z), Cu1(0, 0, 0), Al1(0.06, 0.06, 0), Cu2 (0, 0, z); O1 (0, 0, z), O2(1/2, 0, z), O3(0, 1/2, z); O4(0, 1/2, 0), O5(1/2, 0, 0). Mean square displacements are expressed in Å2. f(O4), f(O5), and f(Al) are the occupational factors of O4, O5, and Al, respectively, defined as atoms per cell.

cold N2 gas blower (Oxford Cryosystems) mounted coaxially. ´ was selected using a doubleA wavelength of λ)0.33483(1) Å crystal Si(111) monochromator. Diffracted intensities were detected through nine scintillator counters, each equipped with a Si(111) analyzer crystal which spans over 16° in the diffraction angle 2θ. The detector bank collected data in the 0 < 2θ < 60° range with a maximum value of the wave-vector Q ) 4π sin θ/λ, Qmax ) 18.7 Å-1 for a total counting time of 3 h. Moreover, data were collected over a wider Q range (Qmax ) 26 Å-1) for one of the samples (x ) 0.33 and δ ) 0.89) using a slightly different wavelength: λ)0.34986(1) Å for a total counting time of 6.5 h. In all cases, data were recorded for longer time at higher angles in order to improve counting statistics. Results The XRPD patterns were analyzed with the Rietveld method as implemented in the GSAS software suite of programs26 which feature the graphical interface EXPGUI.27 The background was fitted by Chebyshev polynomials. Absorption correction was performed through the Lobanov empirical formula28 implemented for the Debye-Scherrer geometry. Line profiles were fitted using a modified pseudo-Voigt function29 accounting for asymmetry correction.30 In the last refinement cycles, scale factor(s), cell parameters, positional coordinates, and isotropic thermal parameters were allowed to vary as well as background and line profile parameters. Deviations from the long-range structure were studied by means of the pair distribution function (PDF), which can be regarded as a suitable technique to map the structure of disordered crystalline phases both in local and intermediate ranges.31,32 In this work the so-called reduced PDF G(r) was used, which indicates the probability of finding an atom at a distance r from another atom and can be obtained via the Fourier sine transform of the experimental total scattering function S(Q)

G(r) )

2 π

∫0Q

max

Q[S(Q) - 1] sin(Qr) dQ

(1)

The G(r) functions related to the experimental data were computed using the program PDFGetX2.33 Data were corrected for diffraction from the capillary, sample self-absorption, and multiple, Laue diffuse and Compton scattering. The analysis of G(r) was carried out with the so-called “Real Space Rietveld” method34 featured by the program PDFGui.35 The degree of accuracy of the analysis is defined by the agreement factor

RW )

n

∑ ω(ri)[Gobs(ri) - Gcalc(ri)]2 i)1

(2)

n

∑

2 ω(ri)Gobs (ri)

i)1

where ω(ri) ) 1/σ2(ri) and σ(ri) is the standard deviation at a distance ri. Data collected up to Qmax ) 26 Å-1 were analyzed starting from r ) 1.6 Å, i.e., including also the shortest Cu-O distances. Data collected up to Qmax ) 18.7 Å-1 were investigated only for values r g 2.6 Å, where robust data reductions were obtained. In fact, these data were affected in the lower r region (r < 2.6 Å) by termination ripples, due to the limited Qmax value. Moreover, disregarding the low r range, changes of the peak widths36 due to thermal motion correlated with the nearest neighbors were avoided.37 The parameters of the long-range structural model obtained by the reciprocal space Rietveld refinement using the diffraction patterns described above are shown in Table 1. The Al-free SmBa2Cu3O6+δ samples are monophasic, while SmBa2Cu2.67Al0.33O6+δ samples are about 98% pure in weight. The impurities are BaAl2O4 (∼1%) and Sm2BaCuO5 (∼1%). The structure of the Al-free sample is tetragonal with space group P4/mmm for

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Figure 2. XRPD pattern collected at 80 K referring to the SmBa2Cu2.67Al0.33O6+δ δ ) 0.89 sample. Observed (crosses) and calculated (continuous line) profiles are shown (top) as well as residuals (bottom). The inset highlights the high-angle region.

Figure 4. Low-temperature conductivity measurements of SmBa2Cu2.67Al0.33O6+δ. The annealing conditions corresponding to each δ value are defined in the text. In the inset: conductivity values (σ) at 273 K for each sample.

Figure 3. High-temperature conductivity measurements of SmBa2Cu2.67Al0.33O6+δ for different pO2: 1 atm (black circles), 10-1 atm (empty circles); 10-2 atm (black squares); 10-3 atm (empty squares).

δ ∼ 0, while it is orthorhombic with space group Pmmm for δ ∼ 1, in accordance with published literature.38 SmBa2Cu2.67Al0.33O6+δ samples exhibit tetragonal structure (space group P4/ mmm) regardless of δ. It should be noted that the atomic fractional coordinates obtained for the Al-free samples, both oxidized and reduced, are in pretty good agreement with the results of Guillaume et al.38 obtained by neutron powder diffraction (using isotopic enrichment). With respect to the Aldoped samples, no neutron diffraction data are available to our knowledge in the literature. However, our previous single crystal X-ray diffraction measurements at room T on SmBa2Cu2.67Al0.33O6+δ samples16 suggested structural models very close to the ones shown in Table 1. The XRPD pattern referring to the (x ) 0.33, δ ) 0.89) sample is shown in Figure 2, as an example. The experimental (crosses) and calculated (continuous line) X-ray diffraction patterns are reported together with the difference curve (bottom). The F(Q) and G(r) functions for the same sample as in Figure 2 are shown in the Supporting Information. The high T conductivity measurements performed on the SmBa2Cu2.67Al0.33O6+δ sample are plotted in Figure 3 versus T. Different symbols are used to tag the associated pO2 values (see figure caption). For pO2 ) 1 and 10-1 atm conductivity decreases monotonically while increasing T. Differently, for pO2 ) 10-2 and 10-3 atm, and for increasing T, conductivity first decreases and then reaches a minimum and eventually increases. The temperature at which the minimum occurs depends on pO2.

The low T measurements performed on “quenched samples” are shown in Figure 4. The δ values reported therein are related to the thermodynamic parameters, T and pO2 of the annealing cycle,15 i.e., T ) 1073 K and pO2 ) 10-4 atm (δ ) 0.47), T ) 1073 K and pO2 ) 10-1 atm (δ ) 0.68); T ) 873 K and pO2 ) 10-1 atm (δ ) 0.82); T ) 873 K and pO2 ) 1 atm (δ ) 0.89); T ) 673 K and pO2 ) 10-1 atm (δ ) 0.94); T ) 673 K and pO2 ) 1 atm (δ ) 0.98). As always happens for this family of superconductors, for samples annealed in the most reducing conditions (i.e., the δ ) 0.47 and 0.68 samples), the resistivity F increases rapidly upon decreasing T. The slope modulus of F(T) decreases while increasing δ and, for δ ) 0.89 and δ ) 0.94 the samples become superconducting at low T. It should be noted that the F ) 0 condition is never reached in the investigated T range, due to the broad transitions to the superconducting state, and TC values were evaluated on the basis of magnetization measurements (see below). As to the δ ) 0.98 sample, F has a metallic behavior for T > ∼150 K, then it shows a broad minimum and, for T < ∼100 K, it increases on decreasing T without superconducting transition. In the whole investigated δ range, however, F decreases monotonically at fixed T against increasing δ. The inset in Figure 4 shows the conductivity (σ) values measured at 273 K as a function of δ. The same data are reported in Table 2, as well as the σ values at 200 K. A stepwise increase of δ is observed for δ g 0.82, i.e., when the superconducting transition takes place. Figure 5 shows the magnetic susceptibility per unit volume subtracted by the high-temperature Pauli contribution from conduction electrons of a selection of the samples represented in Figure 4. The magnetization curves corresponding to the δ ) 0.89 and δ ) 0.94 samples, respectively, are shown in Figure 5 (top). Although only two superconducting samples were investigated, it can already be reasonably inferred that TC depends only weakly on δ. Actually, from Figure 5 one finds that TC ∼ 31.5 K for δ ) 0.89 and TC ∼ 29 K for δ ) 0.94. Very broad transitions to superconducting state are found for



TABLE 2: Conductivity and Mobility Values for “Quenched” Samples at 273 and 200 K, respectivelya

b

δ

σ (273 K) Ω-1 cm-1

σ (200 K) Ω-1 cm-1

0.98 0.94 0.89 0.82 0.68 0.47

110.07 33.72 19.48 6.35 4.67 × 10-1 2.037 × 10-3

134.48 31.31 16.97 4.55 2.12 × 10-1 2.98 × 10-4

∆[h•]Cu2b

µb (273 K), cm2 s-1 V-1

µb (200 K), cm2 s-1 V-1

0.10 0.08 0.055 0.02 0.08

0.606 0.232 0.195 0.175 6.42 × 10-3

0.740 0.215 0.170 0.125 2.91 × 10-3

∆[h•]Cu2c

µc (273 K), cm2 s-1 V-1

µc (200 K), cm2 s-1 V-1

0.129

0.470

0.574

0.066

0.161

0.141

a Mobility values µb and µc are obtained applying the hole injection model of ref 24 and the BVS method described in ref 40, respectively. Hole concentration obtained using the model of ref 24. c Hole concentration obtained using the BVS method described in ref 40.

chamber occurring when air has not completely been vented out; (ii) at T above ∼90 K all data behave in a way which cannot be simply explained by the presence of paramagnetic impurities or Pauli contributions. To clarify this point further investigations with other experimental techniques are needed. Discussion

Figure 5. (top) Field-cooled magnetic susceptibility measured in a field H ) 15 Oe, of two superconducting samples with different O contents and Al ) 0.33 doping. (bottom) Same data for three nonsuperconducting samples with different O contents and Al ) 0.33 doping.

both samples; these can be ascribed to chemical-type inhomogeneities, in a way a natural consequence of the Al for Cu substitution. Conversely, a sharp superconducting transition was stated for the SmBa2Cu3O6.98 sample (see Figure 1 of ref 8). It should be noted that the data were collected in field-cooling conditions at H ) 15 Oe and a reliable estimate of the real Meissner fraction is hence not possible. Indeed, it was shown that an estimate of this fraction in high-TC superconductors is not possible from either zero field cooling or field-cooling measurements, as more complicate calculations are needed.39 In particular, the application of a magnetic field higher than a fraction of Oe (the present case) suppresses most of the diamagnetic Meissner effect. Figure 5 (bottom) displays the magnetic susceptibility curves for the following nonsuperconducting samples: δ ) 0.47, δ ) 0.68, and δ ) 0.98. It should be noticed that the first two samples are “underdoped” while the last is hole-rich (or overdoped). All samples are clearly not superconducting and this is particularly surprising for the overdoped sample. On the other hand, we observe also in all samples: (i) a small bump at T ∼ 40 K, due to the usual O condensation inside the sample

The structural changes induced by oxidation in the structure of YBCO superconductors have been studied by many authors since they are related to hole transfer from the Cu1-O4 chains to the Cu2-O2/O3 planes, where the most mobile charge carriers are localized. In particular, Karppinen and Yamauchi40,41 used the XRPD data obtained by Cava et al.42 on YBCO as a function of δ and found a good correlation between TC and the net hole concentration in the Cu2-O2/O3 planes, calculated through BVM.43 In this method the valence (sij) of the bond between the atomic species i and j is calculated directly from the experimentally determined bond length (dij) by knowing the tabulated bond valence parameters.22,43 The valence is calculated by summing up the bond valences of the i-j bonds to the nearest-neighboring counterions via the so-called bond valence sum (BVS). Following Karppinen and Yamauchi’s40 procedure, a simple BVS could be written considering the net holes residing on copper (Cu2) and oxygen (O2/O3) in the Cu2-O2/O3 planes counteracted by the bonds between Cu2 and apical oxygen (O1) and from in-plane O2/O3 to the Ba and the Sm cation, respectively. For a tetragonal structure, the average plane hole concentration [h•]Cu2, expressed as number of holes per Cu2 ion can be evaluated as follows

[h•]Cu2 ) sCu2-O1 - 4sBa-O2/O3 - 4sSm-O2/O3 + 2

(3) Equation 3 provides a simple way to quantify the increase of the hole concentration in the superconducting planes as a result of structural distortions and fluctuations in atomic potentials.40 Some selected interatomic distances obtained from the Rietveld refinement are reported in the Supporting Information. In Al-free samples, oxidation triggers a noticeable shortening ∆ of the Cu2-O1 bond (∆∼-0.16 Å) as well as an elongation of the Cu1-O1 (∆ ∼ +0.06 Å). At the same time, the Ba atom approaches the z ) 0 plane (∆∼-0.13 Å) without a great increase of the Ba-Cu2 plane distance (∆∼+0.03 Å). In the case of oxidation in SmBa2Cu2.67Al0.33O6+δ compounds only little effect is observed on the Cu2-O1 (∆∼-0.04 Å) bond, whereas the Ba atom approaches the z ) 0 plane (∆∼-0.07 Å). The structural difference between pure and x ) 0.33 doped compound can be related to the decrease of hole transfer to the Cu2-O2 plane due to Al doping, as shown in the discussion about the transport properties.

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Looking at the behavior of bond lengths versus the Al concentration x, it can be stated that the most important structural changes for increasing x at fixed δ () 0.98) are related to Cu1-O1 shortening (∆∼-0.07 Å) and Cu2-O1 elongation (∆∼+0.09 Å). The former effect can be accounted for by the Cu/Al substitution on the Cu1 site, while the latter is consistent with the lower hole transfer rate to the Cu2-O2/3 planes as compared to the Al-free compound. According to the Karppinen method,40 ∆[h•]Cu2 between an oxidized and a reduced (reference) sample is calculated instead of the absolute [h•]Cu2. The reduced sample with δ ) 0.47 was selected as a zero reference for calculating ∆[h•]Cu2. In Table 2 the ∆[h•]Cu2 values for all the investigated samples are reported, in particular one finds ∆[h•]Cu2 ) 0.066 and ∆[h•]Cu2 ) 0.129 holes per Cu2 ion for δ ) 0.89 and δ ) 0.98, respectively. These values will be used in the following while discussing the transport results. Another important structural modification induced by Al doping is the systematic increase of atomic mean square displacements (msd) in doped samples with respect to the Alfree specimen, at the same T (see Table 1). This may be related to local disorder induced by doping: in doped samples msd accounts for the contribution of both thermal atomic vibrations and “static disorder”. As an example, the msd of Ba ions increases from 0.0025 Å2 in SmBa2Cu3O6.98 to 0.0061 Å2 in SmBa2Cu2.67Al0.33O6.98. The same applies to O1 (from 0.004 Å2 in SmBa2Cu3O6.98 to 0.018 Å2 in SmBa2Cu2.67Al0.33O6.98). This is an expected result since Ba and O1 are close to the Cu1 site, where Al substitution takes place. However, a net msd increase is observed even for the Cu2 site (from 0.0009 Å2 in SmBa2Cu3O6.98 to 0.0026 Å2 in SmBa2Cu2.67Al0.33O6.98). In order to prove this hypothesis, real space analysis was performed of the diffraction data for samples with x ) 0.33 and δ ) 0.89 and 0.98, using the Real Space Rietveld method to fit the data, allowing cell parameters, scale factors, atomic positions, and msd’s, i.e., the long-range structural model (space group P4/mmm) found by reciprocal space Rietveld refinement (s. Table 1) to vary over different r ranges. In the following the discussion will be centered on the data collected up to Qmax ) 26 Å-1 from the x ) 0.33 and δ ) 0.89 superconducting sample. However, similar results have been obtained by the analysis of the data collected up to Qmax ) 18.7 Å-1, compatibly with their lower ∆r resolution () 2π/Qmax). Figures and tables pertinent to these last refinements can be found in the Supporting Information. For r > 12 Å the long-range structural model () model_1) performs very well: the fit limited to the bounds 20 Å < r < 50 Å, is shown in Figure 6 “top”. Excellent agreement factors were obtained in the above range. The related structural model parameters, as found from the respective optimizations, are shown in Table 3. At shorter interatomic distances (r < 12 Å) the same model is unsatisfactory. The refinements obtained on the same data in the low r range, produce poor agreement factors (see Figure 6 “middle”). Actually, it should be noted that many local structural features are not well described by the average P4/mmm cell of the long-range structural model, e.g., Al should be surrounded by a tetrahedral environment as determined using MA-NMR for Al-doped YBa2Cu3O6+δ by ref 44 and for Aldoped SmBa2Cu3O6+δ by ref 45. In the above model_1 Al ions exhibit, instead, a distorted octahedral environment with two Al-O1 bonds and four Al-O4 bonds (the O4 sites feature about half occupancy, see Table 1). Furthermore, Al-O and Cu1-O bond distances should be different as a consequence of the different cationic radii of Cu and Al ions.46 Finally, cations such

Scavini et al.

Figure 6. G(r) functions (experimental and best fitted) for the data collected up to Qmax ) 26 A-1 on the SmBa2Cu2.67Al0.33O6.89 sample at 80 K. Observed (crosses) and calculated (continuous line) profiles are shown as well as residuals in the lower part. The width between the dashed horizontal bars of the residuals corresponds to twice the standard deviation (2σ). Top: 20 < r < 50 Å region; the calculated G(r) has been obtained using the crystallographic “average” structural model model_1 (space group P4/mmm). Middle: G(r) functions in the low r range (r < 12 Å); calculated G(r) values are obtained using the same model_1. Bottom: G(r) functions in the low r range (r < 12 Å). The G(r) have been calculated using the supercell model (model_2). The same fits obtained on data collected up to up to Qmax ) 18.7 Å-1 on both δ ) 0.89 and 0.98 samples are shown in the Supporting Information.

Ba and Cu2 with different local environments should occupy different equilibrium positions within the cell. This is testified



TABLE 3: PDF Refinement Results Referring to SmBa2Cu3-xAlxO6+δ δ ) 0.89 Sample at 80 K (Data Collected up to Qmax ) 26 Å-1)a x ) 0.33, δ ) 0.89; Qmax ) 26 Å-1 a ≡ b/Å c/Å Uiso(Sm) z(BaA) z(BaB/zBaC) x,y(BaA) x(BaC) Uiso(Ba) Uiso(Cu1/Al) z(Cu2A) z(Cu2B/Cu2C) Uiso(Cu2) xy(O1A) z(O1A) z(O1B/O1C) Uiso(O1) z(O3) Uiso(O3) Uiso(O4) RW

model_1, 20 < r < 50

model_2, 1.6 < r < 12

3.89011(8) 11.6753(6) 0.00208(10) 0.1856(1)

3.8893(8) 11.6743(6) 0.00114(3) 0.19479(13) 0.18353(4) 0.2441(2) 0.2556(3) 0.00405(6) 0.00506(9) 0.3424(2) 0.35490(5) 0.00123(4) 0.4519(15) 0.1355(8) 0.1642(4) 0.0095(3) 0.3776(1) 0.0087(2) 0.011 0.099

0.25 0.25 0.0069(2) 0.0065(4) 0.3532(2) 0.0041(2) 0.5 0.154(4) 0.036(5) 0.3752(9) 0.0120(13) 0.011 0.084

a

Mean square displacements are expressed in Å2. model_1 uses the same tetragonal P4/mmm long range structural model described in Table 1. Ba, Cu,2 and O1 sites labeled as A, B, and C are equivalent in this model. Concerning model_2, see the text for details. Results obtained on data collected up to Qmax ) 18.7 Å-1 on both δ ) 0.89 and 0.98 sample are shown in the Supporting Information.

by the increase of the msd’s values passing from pure to Aldoped samples at the same T (see Table 1 and Table 3). In order to overcome these problems, the system was modeled using a supercell with sizes 2ap × 2ap × cp metric, where ap and cp are the cell constants of the primitive cell. The supercell considered for the short-range analysis is shown in Figure 7. In particular, in Figure 7 “top” is shown the z ) 0 plane (with Ba ions), while in Figure 7 “bottom” the x ) y plane is displayed. Within this short-range structural model () model_2) the presence of Cu and Al ions in different Cu1 sites has been explicitly considered by placing Al on the (1/2, 1/2, 0) site (actually on the (0.53, 0.53, 0) site, as suggested by ref 16). The remaining Cu sites must have 0.9 Cu and 0.1 Al character in order to obtain the actual sample composition.37 To obtain the tetrahedral coordination for Al in the (0.53, 0.53, 0) site, two O4 ions are placed in (1/2, 3/4, 0) and (3/4, 1/2, 0) and two O vacancies in (1/2, 1/4, 0) and (1/4, 1/2, 0), respectively. The occupational factors of the remaining O4 sites were fixed to 0.39 and 0.48 for δ ) 0.89 and 0.98, respectively, to obtain the correct O concentration. As is evident from Figure 7 “top”, the cell symmetry is decreased while imposing tetrahedral environment for Al, since the 4-fold rotation axis is removed. Applying suitable constraints to the atomic positional degrees of freedom, two symmetry planes have been preserved in the supercell, namely, the x, y z ) 0 and the x ) y, z planes. As a consequence of the symmetry decrease many equivalent atomic positions in the primitive cell are no longer equivalent in the supercell. As an example, the Ba, Cu2, and O1 sites split into three nonequivalent sites each: Ba ions into BaA (1/4, 1/4, z), BaB (3/4, 3/4, z), and BaC (1/4, 3/4, z), Cu2 ions into Cu2A (1/2, 1/2, z), Cu2B (0, 1/2, z), and Cu2C (0, 0, z) and O1 ions into O1A (∼1/2, ∼1/2, z), O1B (0, 1/2, z), and O1C (0, 0, z), as shown in Figure 7.

Figure 7. (Top) z ) 0 plane for a 2 × 2 structural model (Ba ions at z ∼ 0.18-0.20). (Bottom) x ) y plane for the same model. Local distortions highlighted by the refinement are apparent: black circles, Cu ions; black hexagons, Al ions; gray circles, O ions in fully occupied sites; half gray circles, O ions in half occupied sites; empty squares, Ba ions.

A fit considering all the possible degrees of freedom of the supercell would be definitively overparameterized, due to the limited r interval (1.6 < r < 12 Å) considered. In fact, the number of independent data (nd) in the interval ∆r considered in the fit is given by nd ) Qmax∆r/π,32 that is nd ) 86.1 for the Qmax ) 26 Å-1 case. Starting from the structural model obtained in the 20 < r < 50 Å range, many attempts have been made in order to establish the positional degree of freedom which mostly influences the fit without introducing correlations among parameters. Ba, Cu2, and O1 positional disorder seems to affect mostly the fit results. Among others, we tested also a structural model with Al in octahedral environment, without obtaining an improvement of the fit. In model_2, Al-O distances were fixed to 1.80 Å, which is the ideal Al-O distance for Al in tetrahedral coordination.46 Within the BVS assumptions this is equal to supposing that a net charge +3 is present on the Al ion. The only structural degree of freedom allowed to vary for O1A ions is the angle R between the c axis and the Al-O1A direction (see Figure 7 “bottom”). Unlike the case of O1A, zO1B () zO1C) was allowed to vary. The positional degrees of freedom of Ba [zBaA and zBaB ) zBaC, x() y)BaA, xBaC], Cu2 (zCu2A and zCu2B ) zCu2C), and O3 ions (for which all z values are equal) were allowed to vary as well as scale factors, cell constants, and thermal parameters. Thermal parameters associated with equivalent ions in the tetragonal primitive cell are constrained to be all given the same value. A total of 18 parameters have been varied in this model, that is, nd ∼ 4.8 independent data points per parameter. The explicit constraints on the Al-O1 bonds were imposed not only to avoid overparameterization as stated above but also due to the small scattering factors of Al and O ions, when compared to those of Cu, Ba, and Sm. The best fit curve obtained using model_2 is shown in Figure 6 “bottom”. The refinement results are reported in Table 3.

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It turns out that the zBaA coordinate shifts toward higher values, while the z coordinates of the remaining Ba ions decrease with respect to the mean zBa value (see model_1 in Table 3). At the same time, the BaA x ) y position is shifted away from the Al1 ion, while the x coordinate of BaC is displaced toward the fully occupied O4 (1/2, 3/4, 0) site. These rearrangements can be explained by merely electrostatic considerations, taking into account the vacancies introduced by Al doping in the chemical environment of BaA (see Figure 7). For the same reason BaB and BaC are, conversely, attracted toward the z ) 0 plane. Al doping induces disorder also in the positions of “apical” O1 ions, especially in the case of O1A which is directly bound to the dopant Al ion. In particular, O1A shifts toward lower z values (zO1A ) 0.136) and another degree of freedom is allowed along the x ) y direction as a consequence of the tetrahedral coordination of Al imposed by the model. The apical O1 x,y coordinates show a clear shift from the primitive model position, while zO1B ) zO1C ∼ 0.164 is higher than zO1A. It should be noted that similar findings are obtained even removing constraints on Al-O1 distances. Interestingly, disorder seems not to be confined to the neighbors of the Al dopant, but it extends to the Cu2-O3 planes. A net decrease of zCu2A (∼0.342) with respect to zCu2B/C (∼0.355) is actually revealed by the refinement. As a consequence, Al doping causes a corrugation of the superconducting Cu2-O3 planes (Cu2A is 0.15 Å less distant from the Cu1-O4 plane than Cu2B and Cu2C are) as well as a lengthening of the Cu2A-O1A bond (∼2.49 Å) compared to the Cu2B/C-O1B/C bond (∼2.23 Å). Moreover, a nonzero R angle is obtained (R ) 12.4°). The above results deserve some further comments. First of all, we must stress that the model being used is an approximated one: as an example, no positional degrees of freedom are allowed to vary for the Sm ions. Moreover, as stated above, the X-ray coherent scattering cross section of O ions is small if compared with the one of Cu, Ba, and Sm. This implies that subtle differences in metal-oxygen distances have to be handled with care, although the fit is very robust with respect even to O positions. However, with model_2 applied, a neat improvement of the fit quality is obtained in comparison with the aVerage model (model_1). Moreover, many G(r) features at small r values appear to be much better interpreted by model_2 (see Figure 6). The magnitude of the disorder explicitly described through the PDF for Ba and Cu ions is defined as their mean square displacement from their average position (in the P4/mmm cell). This was compared with the square root of the increase of the thermal parameter U observed passing from SmBa2Cu3O6.98 to SmBa2Cu2.67Al0.33O6.89 (see Table 1). In the case of Ba ions, a very good agreement was found: 0.056(2) Å in the reciprocal space compared to 0.057(2) Å obtained according to the supercell model in the real space (see Table 1). At the same time, the residual thermal parameter in the real space diminished from ∼0.007 to ∼0.004 Å2 (see Table 3). Even in the case of Cu2 ion, a good matching both in real and reciprocal space has been obtained. In particular, real space root mean square deviations were found to be slightly bigger (∼0.054(5) Å) than the ones in the reciprocal space (∼0.040(3) Å). Even in this case the isotropic thermal parameter U decreases from ∼0.004 to ∼0.001 Å2 (see Table 3).

Scavini et al.

Figure 8. High-temperature conductivity measurements of SmBa2Cu2.67Al0.33O6+δ as a function of δ. Different symbols refer to different annealing pO2: 1 atm (black circles); 10-1 atm (empty circles); 10-2 atm (black squares); 10-3 atm (empty squares).

The same conductivity data reported in Figure 3 are depicted in Figure 8 in a logarithmic scale as a function of δ. The same symbols as in Figure 3 are used to label the data according to the pO2 values at which the data were collected. The δ values are taken from ref 15. While varying δ, three different conductivity regimes can be identified. For δ > ∼0.78 (zone I in Figure 8) σ increases by increasing δ. No or weak dependence of σ on the annealing pO2 is apparent. Since the same δ value is obtained by varying pO2 at different T,15 in this δ-range the variation of the carrier mobility µ as a function of T is immaterial. Also in the “intermediate” δ range 0.6 > ∼δ > ∼0.78 (zone II in Figure 8) σ increases upon increasing δ. In this range, however, σ(δ) clearly depends on the annealing pO2. As at increasing pO2 the same δ value is obtained upon increasing T, it is possible to conclude that in zone II thermally activated conductivity is present. Unlike zones I and II, in the “O-depleted” zone III (δ < ∼0.6) σ increases versus decreasing δ. A similar sign change of the slope of the σ(δ) curve had been assessed also for the Al-free SmBa2Cu3O6+δ material23 but at a lower δ value (δ < ∼0.20). In the case of ref 23, conduction in the dσ/dδ < 0 zone was attributed to the intrinsic ionization giving rise to couples of electronic defects, since holes introduced by oxidation in this δ range should be strongly localized on Cu1 sites. The presence of an energy gap at low δ (ideally δ ) 0) for YBCO was inferred by Maier and Tuller.47 Differently from them, Chiodelli et al. attributed charge transport in this zone to ionic conductivity.48 The difference in the δ minimum of the conductivity between SmBa2Cu2.67Al0.33O6+δ and SmBa2Cu3O6+δ compounds (∆δ ∼ 0.40) can be at least partially accounted for by the equilibria of the active defects in the two compounds.15 Actually, if SmBa2Cu3O6 is chosen as the reference structure, for each substituting Al ion two holes are annihilated.15 Further information about the transport properties of SmBa2Cu2.67Al0.33O6+δ can be gained by taking into account the σ measurements at low T as shown in Figure 4. As previously stated, for samples with O low nonstoichiometry δ values (i.e., δ ) 0.47 and 0.68), which in Figure 8 correspond to zones III and II, respectively, F increases rapidly versus decreasing T and the slope modulus of F(T) decreases when δ increases. Samples with higher δ (zone I of Figure 8) exhibit either a lower F(T) slope (δ ) 0.82), or a supercon-

Structural Disorder Induced by Al Doping ducting transition (δ ) 0.89, 0.94) or, further on, metallic conduction (δ ) 0.98). For each T, conductivity increases with increasing δ. If conductivity at a fixed T is plotted against δ, a noticeable increase of σ happens only for δ g 0.82 (see inset of Figure 4 for measurements at 273 K and Table 2). Attention should be drawn to the fact that the analysis of the transport properties of this compound would be more informative if the study of the charge carrier conductivity were substituted by the investigation of µ, since this parameter is more strongly related to the transport mechanism. Unfortunately, varying δ, the hole injection sequence is quite complex in the superconductors of the YBCO family24,49 making it difficult to convert σ data into µ values. As pointed out by Tolentino et al. who performed XAS measurements on YBa2Cu3O6+δ samples,49 holes introduced by O doping are localized on Cu1 and on the O sites. Only a fraction of holes are transferred to the superconducting Cu2-O2/3 planes. The latter are the most mobile charge carriers and are responsible for the superconductivity mechanism.49 The conductivity results obtained at both low and high T can be interpreted considering the simplified model for hole injection as a function of δ, as proposed by Spinolo et al. for SmBa2Cu3O6+δ.24 Within this model holes injected into the structure at low δ are either largely localized on the Cu1 sites (δ < 0.20) or have strong O4-2px,y character (0.20 < δ < 0.50) and scarce mobility. Above δ ) 0.50 holes have O2/3-2px,y character and are produced at a rate of 1 hole/doping O. These holes are the most mobile, behave like large polarons, and are responsible for superconductivity.24 By deploying the model in ref 24 for the Al-doped compound, one can say that holes introduced into the structure at low δ values (zones II and III in Figure 8) remain confined in the Cu1-O4 plane and are rather localized while mobile holes are injected into the superconducting Cu2-O2 plane only for high δ values, i.e., δ > ∼0.50 for SmBa2Cu3O6+δ and δ > ∼0.78 (zone I) for SmBa2Cu2.67Al0.33O6+δ. When pure and Al-doped SmBa2Cu3O6+δ are compared (cf. Figure 1 in ref 24 and Figure 6 in ref 23) some differences in the transport properties are apparent. In particular, in Al-doped samples: (a) the rise of TC is shifted to higher δ values (δ ) 0.89 instead of δ ) 0.50); (b) higher F values are obtained, compared to the Al-free compound; (c) dF/dT < 0 for almost all the Al-doped samples, even for the superconducting specimens; (d) the only exception to (c) is the δ ) 0.98 sample for T > 100 K. This sample is not superconducting (see Figures 4 and 5), and should be assigned to the overdoped region. While (a) can be (at least partially) accounted for by the reduction of the total hole concentration introduced by Al doping, (b-d) needs to be better analyzed. The conductivity drop can derive from a decrease of charge carrier concentration and/or mobility. In the former case only defect equilibria have to be taken into account, while in the latter also the transport mechanism is involved. To this purpose a rough estimation of mobility at a fixed T (e.g., at 273 K, see the inset of Figure 4) can be given using the method suggested in ref 24 supposing that (i) holes introduced into zone III (δ < 0.60) are fully localized on Cu1 sites and must not be considered, (ii) holes introduced into zone II (0.60 < δ < 0.78) have mainly O4-2px,y character, while (iii) holes introduced above δ ) 0.78 (zone I) feature O2/3-2px,y character. In zones I and II, holes are produced at a rate of 1 hole/doping O. The charge carrier concentration values, expressed as holes/Cu site, are reported in Table 2. In the same table the values obtained by BVS are also shown for δ ) 0.89 and 0.98. The two methods


Figure 9. Arrhenius plot of the resistivity data shown in Figure 4. The inset shows the apparent activation energy for conduction.

give similar results, the discrepancies being attributed either to a slightly incorrect determination of the boundary between zone I and II or to the oversimplified model adopted for hole transfer to the superconducting planes or to both of them. The mobility values obtained at 273 and 200 K using the two methods are shown in Table 2. (Mobility has been evaluated by applying the formula σ ) [h•]Cu1qµCu1 + [h•]Cu2qµCu2, where [h•] is the charge carrier concentration. The subscripts Cu1 and Cu2 refer to the plane where the carriers are confined. Since µCu1 , µCu2, in zone I the above equation can be simplified into: σ ≈ [h•]Cu2qµCu2.) The only sample pertaining to zone II is δ ) 0.68. For this sample µ ∼ 0.006 cm2 s-1 V-1 at 273K and µ ∼ 0.003 cm2 s-1 V-1 at 200 K. When zone I is reached (δ ) 0.82, 0.89, 0.94) at 273 K, mobility increases by nearly 2 orders of magnitude (µ ∼ 0.2 cm2 s-1 V-1). In this δ range µ increases smoothly at increasing δ. It should be noted that when the BVM is used to determine the charge carrier concentration and/or lower T are considered, smaller mobility values are obtained (see Table 2). Only for the most oxidized sample (δ ) 0.98) µ ∼ 0.6 cm2 s-1 V-1 at 273 K. The large increase of charge carrier mobility stepping from zone II to zone I is consistent with the hypothesis that only in zone I (for δ ∼ >0.78) mobile are holes localized in the Cu2-O3 planes. However, mobility values for SmBa2Cu3O6+δ superconducting samples (i.e., for δ > 0.5) are always well above 1 cm2 s-1 V-1; the same holds for the Nd-123 isostructural compound.50 These mobility values are almost 1 order of magnitude higher than our experimental findings; thus the decrease of conductivity detected for Al-doped samples compared to the Al-free reference compound has to be related to some variation of the transport mechanism triggered by Al doping. With regard to point c above, dF/dT < 0 is usually observed in the tetragonal nonsuperconducting low-δ zone of LnBa2Cu3O6+δ materials. This behavior is deemed to be due to the presence of Anderson localization, namely, fluctuations in the band structure near the Fermi level, driven by structural disorder.24,50,51 As opposed to that, superconducting samples display “metallic”-like conductivity (dF/dT > 0).24,50 Mobile holes in the Cu2-O2/3 planes have been meant to behave like large polarons in the case of both Al-free Sm-12324 and isostructural Nd-123 compound.50 As stated before, for almost all Al-doped Sm-123 samples (except for δ ) 0.98) dF/dT < 0; this is true even for superconducting samples far from the superconducting transition. In Figure 9 the Arrhenius plot (log σ vs 1/T) of the low-T conductivity data from Figure 4 are shown for selected samples.

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Figure 10. Logarithm of the conductivity σ versus T-1/3 plot for the δ ) 0.82 sample.

The high-T part of the data has been used to determine the “apparent” activation energy E for conduction. The inset shows the E values determined for all the SmBa2Cu2.67Al0.33O6+δ samples as a function of δ. E decreases monotonically when δ increases. E ) 3.4 kcal/mol in zone III (δ ) 0.47), and E ) 1.7 kcal/mol in zone II (δ ) 0.68). Within zone I, E varies from 0.36 kcal/mol for δ ) 0.82 to 0.20 and 0.19 kcal/mol for the two superconducting samples (δ ) 0.89, 0.94), respectively. E has hence a finite value even for the superconducting samples; only the most oxidized sample (δ ) 0.98) is characterized by a metallic behavior, at least for T > 100 K. Focusing on zone I, it is possible to conclude that on increasing δ (and hole concentration bounded in the Cu2-O3 planes), E decreases monotonically and eventually vanishes. At the same time, µ at a fixed T (e.g., at 273 K, see Table 2) increases from ∼0.18 to ∼0.61 cm2 s-1 V-1. The mobility difference increases when lower T are considered (e.g., at 200 K, see Table 2). All these features are consistent with the presence of Anderson localization. In fact, although Al ions substitute Cu only in the Cu1 site, i.e., far from the superconducting planes, disorder spreads even over the Cu2-O3 planes, as shown previously. In particular, model_2 provides a net decrease of the Cu2A z coordinate (see Table 3). As a consequence, the “semiconducting”-like low-T conductivity behavior of the samples belonging to zone I could be related to a variable range hopping (VRH) mechanism due to energy fluctuations induced by structural disorder. Figure 10 shows the log σ vs T-1/3 plot for the δ ) 0.82 sample. The linearity of the low-T data (for 25 < T < 110 K) is the fingerprint of the Mott-Davis bidimensional VRH mechanism.52 Accordingly, the “apparent” E value for charge transport decreases monotonically while increasing the hole concentration. The coexistence of superconductivity and Anderson localization has been observed in other superconducting compounds such as Nd2-xCexCuO4+δ.53 In that case structural disorder is introduced by the shrinkage of the O cuboid around the Ce dopant along the crystallographic c axis, which, in turn, causes the corrugation of the superconducting CuO4 planes.54 As to SmBa2Cu2.67Al0.33O6.98, this compound shows “metallic”-like conductivity for ∼150 < T < 300 K; conversely, the related dF/dT is less than zero for lower T, and it never undergoes a transition to the superconducting state, as shown by both conductivity and magnetization measurements (see Figures 4 and 5). In addition, µ ∼ 0.6 cm2 s-1 V-1 at 273 K, i.e., the mobility is of the same order of magnitude as that of the Al-free SmBa2Cu3O6+δ compounds.24 The increase of

Scavini et al. mobility and the crossover between metallic and actiVated conduction suggest that the charge carrier concentration for this δ is next to the mobility edge of Anderson localization. The absence of superconducting transition indicates that this sample belongs to the so-called overdoped region. Generally, the suppression of superconductivity within the overdoped region takes place at much higher carrier concentrations (0.25-0.30 holes per Cu ion; see, e.g., ref 55). Suitable doping on the Cu sites can however decrease the [h•]Cu2-range of the TC([h•]Cu2) parabolic curve.1,56 The above discussion emphasizes that the changes in transport properties induced by Al-doping in SmBa2Cu3-xAlxO6+δ are not explained by the mere decrease of hole concentration at fixed δ, due to the formation of AlCu1•• defects (using SmBa2Cu3O6 as a reference structure15). As testified by the real space analysis of XRPD patterns, disorder induced in the superconducting Cu2-O3 planes causes Anderson localization to occur even for holes lying in the same planes. Moreover, the mobility edge is shifted toward the overdoped zone and the p range of TC([h•]Cu2) is strongly reduced. The presence of this kind of disorder could help to shed light on some nontrivial effects occurring in SmBa2Cu3-xAlxO6+δ samples such as the increase of TC and Meissner fraction by varying the Al distribution within the Cu1 sites while keeping x and δ constant (i.e., the total charge carrier concentration19 in SmBa2Cu2.85Al0.15O6.98), as well as the appearance of superconducting diamagnetic fluctuations above TC in the same compound. As stated previously, Al doping promotes corrugations within the superconducting Cu2-O3 planes. In particular, the supercell model used in the PDF analysis showed that Cu2A is 0.15 Å less distant from the Cu1-O4 plane than Cu2B and Cu2C are (see above). This is equivalent to arguing that potential fluctuations occur in the superconducting planes. At fixed x and δ, not only the Al concentration but also its distribution within the Cu1 sites can thereby influence the charge carrier diffusion path within the Cu2-O3 planes. When Al clustering is induced in the z ) 0 plane, Al-free and Al-rich zones are induced within the same plane. In previous papers19,20 we have shown that the size of Al-free zones is of the order of 20 Å in SmBa2Cu2.85Al0.15O7. As can be concluded from our present results, potential fluctuations with equal spatial extension should be induced also in the Cu2-O3 planes, affecting the superconducting properties, since the coherence length along the ab direction of superconducting carriers in RE123 superconductors was found to be close to this value.57 For what concerns REBa2Cu3O6+δ superconductors, Al doping and clustering seems thus to be a suitable method to investigate the intriguing phenomena which appear in the “underdoped zone” of these superconductors. In fact, it is possible to produce samples with the desired charge carriers concentration (tuning x and δ) and spatial extension of potential fluctuations within the Cu2-O2/3 planes using appropriate annealing cycles.17,19 Summary In the present study we investigated the effect of local disorder on the transport properties of Al-doped SmBa2Cu2.67Al0.33O6+δ superconductors. High-resolution XRPD measurements revealed that for fully oxidized samples (δ ∼ 1) Al-doping causes an elongation of the Cu2-O1 bond, as well as a shortening of the Cu1-O1 bond. Applying the BVS method, it was shown that this causes a decrease of hole transfer to the Cu2-O3 superconducting planes. Moreover, Al-doping leads to the systematic increase of thermal parameters at fixed T. The real space analysis of the

Structural Disorder Induced by Al Doping diffraction data using the PDF technique revealed that Al doping introduces local disorder in the proximity of the dopant Al ion, especially in the Ba, O1, and Cu2 positions. Disorder spreads to the Cu2-O3 planes producing corrugations in these planes. Conductivity and magnetization measurements proved that superconductivity is attained only for δ g 0.89. The compound with δ ) 0.98 belongs to the overdoped region and, for it, TC ) 0. Charge carriers are transferred to the Cu2-O3 planes only for δ ∼ >0.78 and mobility increases of about 2 orders of magnitude. However, lower mobility values are obtained for doped samples, as compared to the Al-free reference material(s) and “activated” mobility is attained even for superconducting compositions. This behavior has been ascribed to the appearance of Anderson localization induced by the structural disorder and potential fluctuations present in the Cu2-O3 planes as a direct consequence of Al doping. Since it is possible to control the charge carrier concentration (tuning x and δ) and the spatial extension of potential fluctuations within the Cu2-O2/3 planes using appropriate annealing cycles, Al doping and clustering seems to be a suitable method to investigate the intriguing phenomena which appear in the “underdoped zone” of REBa2Cu3O6+δ superconductors. Acknowledgment. The authors gratefully acknowledge the European Synchrotron Radiation Facility for provision of beam time. They are also greatly indebted to Dr. Andy Fitch for assistance in using the ID31 beamline. They would like to thank the ISTM-CNR Institute in Pavia, especially Dr. Gaetano Chiodelli for the kind hospitality and the use of their instrumentation for high- and low-temperature conductivity measurements. The EU-Network of Excellence MAGMANet is acknowledged for partly funding the magnetic measurements. Supporting Information Available: Tables reporting selected interatomic distances (in Å) obtained by the structural models described in Table 1 and the PDF Refinement results referring to SmBa2Cu3-xAlxO6+δ δ ) 0.89 and 0.98 sample at 80 K (data collected up to up to Qmax ) 18.7 Å-1) and figures reporting the F(Q) and G(r) functions for the SmBa2Cu2.67Al0.33O6+δ δ ) 0.89 sample collected at 80 K and the G(r) functions (experimental and best fitted) for the data collected up to Qmax ) 18.7 A-1 on the SmBa2Cu2.67Al0.33O6.89 and SmBa2Cu2.67Al0.33O6.98 samples. This material is available free of charge via the Internet at http://pubs.acs.org. References and Notes (1) Naqib, S. H.; Cooper, J. R.; Tallon, J. T.; Panagopoulos, C. Physica C 2003, 387, 365–372. (2) Hinkov, V.; Bourges, P.; Pailhe´, S.; Sidis, Y.; Ivanov, A.; Frost, C. D.; Perring, T. G.; Lin, C. T.; Chen, D. P.; Keimer, B. Nat. Phys. 2007, 3, 780–785. (3) Koblischka, M. R.; Winter, M.; Das, P.; Koblischka-Veneva, A.; Muralidhar, M.; Wolf, T.; Hari Babu, N.; Turner, S.; van Tendeloo, G.; Hartmann, U. Physica C 2003, 469, 168–176. (4) Larkin, A. I.; Varlamov, A. A. Theory of Fluctuations in Superconductors; Oxford University Press: Oxford, 2005. (5) Carretta, P.; Lascialfari, A.; Rigamonti, A.; Rosso, A.; Varlamov, A. A. Phys. ReV. B 2000, 61, 12420. (6) Lascialfari, A.; Rigamonti, A.; Romano`, L.; Tedesco, P.; Varlamov, A.; Embriaco, D. Phys. ReV. B 2002, 65, 144523. (7) Lascialfari, A.; Rigamonti, A.; Romano`, L.; Varlamov, A. A.; Zucca, I. Phys. ReV. B 2003, 68, 100505 (R). (8) Bernardi, E.; Lascialfari, A.; Rigamonti, A.; Romano`, L.; Scavini, M.; Oliva, C. Phys. ReV. B 2010, 81, 064502. (9) Andersen, N. H.; von Zimmermann, M.; Frello, T.; Ka¨ll, M.; Mønster, D.; Lindgård, P.-A.; Madsen, J.; Niemo¨ller, T.; Poulsen, H. F.; Schmidt, O.; Schneider, J. R.; Wolf, Th.; Dosanjh, P.; Liang, R.; Hardy, W. N. Physica C 1999, 317-318, 259–269.

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JP106805Z

Effect of Local Disorder on the Transport Properties of Al-Doped SmBa

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