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Characterisation of La2NiO4+δ using in-situ high temperature neutron powder diffraction
Solid State Sciences 5 (2003) 419–426 www.elsevier.com/locate/ssscie
Characterisation of La2NiO4+δ using in-situ high temperature neutron powder diffraction Stephen J. Skinner Centre for Ion Conducting Membranes, Department of Materials, Imperial College London, Prince Consort Road, London, SW7 2BP, UK Received 14 October 2002; received in revised form 16 December 2002; accepted 8 January 2003
Abstract La2 NiO4+δ , has been studied using in-situ high temperature neutron diffraction over a temperature range of 25–800 ◦ C in vacuum. The behaviour of this material, and in particular the oxygen interstitial content, is discussed and quite remarkable bond length changes observed. It is observed that at temperatures above 150 ◦ C La2 NiO4+δ transforms to the tetragonal I 4/mmm structure and maintains this over the entire temperature range on both heating and cooling. The loss of the interstitial oxygen was observed over the low temperature region of the study and significant changes in both lattice constant and bond lengths found to mirror these changes, indicating the structural importance of the interstitial oxygen. 2003 Éditions scientifiques et médicales Elsevier SAS. All rights reserved. Keywords: In-situ; Neutron diffraction; High temperature; Interstitial; K2 NiF4 ; La2 NiO4+δ
1. Introduction Although the structure of La2 NiO4+δ has been extensively investigated [1–7] there have been no in-situ high temperature studies performed, either by XRD or neutron diffraction. It has been reported that materials of the K2 NiF4 structure type can accommodate a wide variety of oxygen stoichiometries, predominantly hyperstoichiometry. This has been found to influence the symmetry of La2 NiO4+δ , with a variety of structural modifications based on the tetragonal I 4/mmm spacegroup reported. La2 NiO4+δ adopts the K2 NiF4 structure but with a slight orthorhombic distortion of the unit cell in the ab plane [6]. The stoichiometric material can be easily oxidised with the excess oxygen occupying interstitial sites and, depending on the amount of oxygen incorporated, the structure will vary between related tetragonal and orthorhombic structures. Several studies, including some utilising neutron diffraction techniques [5–8], have accurately determined the degree of orthorhombic distortion and the position of the oxygen interstitials. However there remains considerable debate as to the exact nature of the low temperature form of La2 NiO4 with both orthorhombic and tetragonal unit cells having been proposed, as well as biphasic mixtures. Indeed it has been proE-mail address: [email protected] (S.J. Skinner).
posed that a two phase region of the phase diagram exists for compositions with 0.02 δ 0.13 [6] and further details of oxygen intercalated samples indicate three two phase regions existing between 0.03 δ 0.13 [3]. The Fmmm phase has been proposed as being stable above the δ = 0.13 limit. Synthesis method has also been suggested as an influencing factor in determining the phase produced [7]. Unfortunately to date there has been little investigation of the behaviour of these oxides in-situ at elevated temperatures. This may partly be due to the lack of interest in high temperature applications of K2 NiF4 type oxides as the majority of efforts in recent years have been directed towards the La2 CuO4 related superconducting oxides and hence most studies have concentrated on the materials at < 25 ◦ C. One study, of La2 NiO4.13 , has been reported [9] utilising neutron diffraction techniques but the details of the report are limited to a maximum temperature of 400 ◦ C. A subsequent study by Paulus et al. [8] details structural refinements on single crystal samples at three temperatures, below 25 ◦ C, at 25 ◦ C and at 400 ◦ C where the materials have been electrochemically oxidised. A further study [10] has endeavoured to use heat capacity anomalies in a single crystal to identify structural changes, but as previously the study was limited, with a maximum temperature of 227 ◦ C. However, there has been increasing interest in K2 NiF4 type oxides as possible oxide ion conducting ceramics since the discovery
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of mixed ionic-electronic conduction in the La2 NiO4+δ material [11–13]. Ion conducting ceramics are widely used in high operating temperature devices such as solid oxide fuel cells and ceramic oxygen generators. It is therefore essential to understand the stability and properties of these oxides at elevated temperatures if La2 NiO4+δ and related materials are to be considered as new mixed ionic-electronic conducting components in electrochemical devices such as solid oxide fuel cells. In this paper the phase stability and transitions of La2 NiO4+δ is examined and the behaviour on both heating and cooling in vacuum investigated.
2. Experimental La2 NiO4+δ was prepared using a modified Pechini [14] method in which solutions of the appropriate nitrates were combined with excess citric acid. This mixture was then heated until gelation occurred and then dehydrated and decomposed at 500 ◦ C. After decomposition the material was sintered at 1370 ◦ C for 12 h after which a single phase, as determined by XRD, was identified. In order to establish the oxygen stoichiometry of the phase produced thermogravimetric analysis in a 10% H2 /N2 mixture was performed using a Stanton Redcroft 780 series instrument. The reduction of the oxygen excess material to the component oxides and metals allowed the accurate determination of the value of δ. The in-situ neutron diffraction measurements were carried out on the HRPD beamline at ISIS at the Rutherford Appleton Laboratories. Measurements were carried out in the vacuum conditions of the apparatus to enable a monitoring of the loss of excess oxygen as temperature increased. Further, to enable any structural modifications resulting from oxygen loss to be detected both heating and cooling cycles were recorded. Heating of the samples was carried out using an induction furnace which would easily achieve the 800 ◦ C upper temperature required.
3. Results and discussion Considerable debate surrounds the crystal structure adopted by La2 NiO4+δ both at room temperature and at elevated temperatures [2,3,5–8,15–20] and in order to elucidate the structure at elevated temperature the high resolution powder diffraction (HRPD) instrument was used to investigate these phases. Data was recorded at temperatures of up to 800 ◦ C, a temperature which was deemed to be sufficient to allow the evolution of the interstitial oxygen without causing the reduction of the material to the component oxides and metals. To determine the oxygen stoichiometry thermogravimetric analysis was carried out under the previously described reducing conditions and an oxygen excess content of
0.18 ± 0.01 was determined for the as synthesised material at room temperature. Initial neutron diffraction data recorded at room temperature and in vacuum indicated that a small amount of NiO was present in the sample (∼ 2%) which was not detected in the powder XRD patterns. Due to the relatively strong scattering of the Ni ion the small amount of NiO detected is quite prominent in the neutron powder diffraction pattern. The peaks observed that were attributable to NiO were clearly distinguishable from those due to the La2 NiO4+δ phase. Obviously this observation would have an effect on the TGA results, but as the NiO would not reduce until the same temperature as the La2 NiO4 the initial weight loss at ∼ 400 ◦ C would be attributable to the interstitial loss. Therefore the small percentage of NiO present in the sample has a minimal effect on the determination of the oxygen nonstoichiometry. Subsequent structural refinements using the GSAS software [21] were therefore two component models including the Fm3m cubic NiO phase with lattice parameters of a ∼ 4.18 Å. In the initial refinements the La:Ni ratio was refined but found to be as the starting stoichiometry suggested and therefore in later refinements the occupancy of these sites was fixed at 1. The presence of the NiO phase did have a benefit to the refinement process in that the lattice parameter expansion of the oxide phase provided a useful guide in terms of the thermal expansion effect on the lattice of both materials. A completely linear behaviour on both heating and cooling was observed which indicated that the data collection processes were reliable and any changes observed in the La2 NiO4+δ diffraction patterns would be due to changes within the material and not instrumental. Several authors have proposed models for the structure of La2 NiO4+δ at room temperature and include both single phase and two phase models depending on the oxygen stoichiometry. Therefore Rietveld refinements of several possible models were initially attempted using the GSAS suite of programmes [21] and including both tetragonal and orthorhombic possibilities. A shifted Cherbyshev background function and pseudo Voigt peak shape were used in the refinement process. Early in this process two phase models were discounted as there was no evidence to support this from the experimental data through either peak shape, broadening or splitting. Contrary to the findings of Paulus [9] and Mehta [2], models fitting to space groups Bbcm and F 4/mmm were found to give unrealistic fits to the experimental data. Furthermore, attempts to refine the data to the tetragonal I 4/mmm space group resulted in a physically implausible model. A best fit to the room temperature data was obtained in space group Fmmm with the La occupying the (00z) position, Ni occupying the (000) position and oxygen occupying four positions with variable fractional occupancies, as detailed in Fig. 1 and Table 1. The refined data also enabled the oxygen content to be determined with some accuracy and it was found to be in good agreement with the TGA data, giving a final refined value over all oxygen positions of δ = 0.17. The
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421
Fig. 1. Neutron diffraction time-of-flight data for La2 NiO4+δ recorded at 25 ◦ C. (+) represent experimental data points, solid line is refined data fit. The upper markers are the minority NiO phase, lower markers are La2 NiO4+δ orthorhombic phase. Rp = 3.92%, Rwp = 4.57% (d-spacing range 0.6–2.5 Å). Table 1 Refined atomic positions for La2 NiO4+δ in Fmmm spacegroup. Data recorded at 25 ◦ C Name La1 Ni1 O1 O2 O3 O4
x
y
z
0 0 0.25 0.25 0 −0.0661(17)
0 0 0.25 0.25 0 −0.0627(16)
0.36113(7) 0 0 0.22674(167) 0.17421(27) 0.16990(50)
Occupancy 1 1 1.0031 0.038(4) 0.531(20) 0.119(6)
occupancy of the 4 oxygen positions is also in reasonable agreement with the work of Jorgensen et al. [6] who refined the room temperature structure of La2 NiO4+δ in the Fmmm spacegroup. The standard deviation for the oxygen stoichiometry is subject to a reasonably high degree of uncertainty due to the fitting of the O2 position in particular as the error in the site occupancy is of the order of 10% which is related to the low occupancy of this site, 0.038. It is also interesting to note that the O2 position has significant variation in the positional parameters as well as the thermal parameters. This uncertainty is also likely to be due to the virtually zero value for the site occupancy. This contrasts
U11 0.81(2) 0.75(2) 1.31(14) 0.2 0.21(14) 0.67(22)
U22 0.81(2) 0.75(2) 0.50(13) 0.2 0.21(14) 0.67(22)
U33 0.81(2) 0.75(2) 2.69(9) 0.2 0.21(14) 0.67(22)
U12 0 0 −0.06(5) 0 0 0
with the accuracy of the determination of the position of the fully occupied site, O1, which has a standard deviation of 0.08%. Hence, it is indicative of the oxygen stoichiometry of the sample and when used in conjunction with the data obtained from the TGA measurements can confirm the total stoichiometry of the sample and the overall oxygen losses from the material. Consequently the structure of this sample at room temperature is found to be orthorhombic Fmmm with refined lattice parameters of a = 5.45869(4) Å, b = 5.46360(4) Å and c = 12.68524(4) Å. With the room temperature data refined and in good agreement with the literature the main focus of the work
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was on the elevated temperature data sets. Two sets of lower temperature data were collected, at 150 and 300 ◦ C, with the data then collected every 50 ◦ C from 400 until 700 ◦ C. A final set of data was recorded at 800 ◦ C and then every 100 ◦ C on cooling until 300 ◦ C. Initial analysis of the lattice parameter data on heating indicated that there was non-linear behaviour in the c parameter and very little change in the a and b parameters. The interpretation of this data was assumed to be due to the oxygen loss from the La–O layers. The temperature at which this transition was evident was between 300 and 700 ◦ C and this temperature range corresponded well to the data collected from thermogravimetric analysis. Rietveld refinements of the experimental data at each temperature were performed using GSAS and each data set carefully examined for evidence of tetragonal/orthorhombic distortions. From the data recorded at 150 ◦ C, Fig. 2, there was no evidence of orthorhombic splitting and/or line broadening indicative of an orthorhombic distortion. The model used in this refinement was based on the I 4/mmm space group with an oxygen excess position at the (0 0.5 0.23) position. The refinement details are given in Table 2. In this model the anisotropic temperature factors were refined only for O2 as the refinement of this position isotropically did not give a physically plausible model. Refinement of all other temperature factors using an isotropic model, as observed in Table 2, resulted in temperature factors with reasonable values. The values for O2 gave a particularly high Uiso and on subsequent analysis of the anisotropic temperature factors it is evident that this was due to the large difference between the U11 /U22 values and that for U33 . The large ab plane anisotropy of the apical oxygen position in K2 NiF4 type oxides is well known and indeed may contribute to the fast oxide ion conduction in these materials, and is illustrated by the fitting of the temperature factors of the O2 position. To ensure that the spacegroup assignment was correct data refinements were performed in the orthorhombic spacegroups detailed previously but there were no significant improvements in the fits or indeed in the Rwp and Rp values. From the refinement in the tetragonal spacegroup it was possible to determine an oxygen excess stoichiometry of δ ≈ 0.13. Further refinements up to 800 ◦ C revealed no further structural transformations and all data could be refined in the I 4/mmm space group. Typical data from a refinement at 600 ◦ C are given in Table 3. Only the change in oxygen stoichiometry was evident when the lattice parameters were examined and the oxygen positions refined. As anticipated from previous studies the overall oxygen excess concentration determined from the Rietveld refinement of the oxygen positions decreased from ∼ 4.17 at room temperature to ∼ 4.00 at 450 ◦ C. Examination of the variation in the lattice parameter data with temperature, Fig. 3, is revealing. On heating there is in fact very little change in both the a and b parameters (orthorhombic lattice parameters from the room temperature
√ data have been converted through a 2a relationship to the tetragonal cell) but on studying the behaviour of the c parameter it is obvious that there is a large change associated with the loss of oxygen. It is also telling that at both the low and high temperature regions the c parameter behaviour is linear, whereas there are significant differences in the behaviour around the transition temperature range. As the oxygen stoichiometry has undergone significant change at the lower temperature it is interesting to investigate the reason for the linear behaviour of the c parameter in this region. In the transition region there are two main competing demands on the structure—adjustment to the loss of interstitial oxygen and thermal expansion. The linearity of the expansion in the c cell constant at temperatures below 400 ◦ C and above 650 ◦ C can be attributed to thermal expansion whereas between 400 and 650 ◦ C there is the additional loss of oxygen modifying the structure. These competing demands explain the deviation from linearity in the lattice parameters with temperature and also explain the linearity of the plot on cooling, where there is no variation in oxygen stoichiometry. It is also well known [2,6] that the oxygen interstitial site lies in the La–O plane in the ab direction close to the (0.25, 0.25, 0.25) position and therefore further evidence for the loss of oxygen from interstitial sites should be evident from the bond distance data. Furthermore, these reports indicate that the incorporation of the interstitial oxygen forces the apical oxygen to be forced into two split positions (0, 0, ∼ 0.17) and (−0.06, −0.06, 0.17) and the resultant effect on bond distance should be evident. Looking at the detail of the bond distances it is of interest to note the behaviour of the equatorial and apical Ni–O bonds in the NiO6 octahedron. It was anticipated that the Ni–O bonds would lengthen with increase in temperature and indeed this was observed for both the equatorial and apical positions, Fig. 4. However, on closer examination it is evident that the two apical Ni–O bonds have experienced a much greater effect than the equatorial positions. It is also interesting to observe that the apical position bond lengths undergo a dramatic lengthening at ∼ 400–450 ◦ C, indicating that this change was directly attributable to the loss of interstitial oxygen ions. It is likely that in the oxygen excess materials there is a small quantity of Ni3+ present. If this is the case then the Ni3+ –O bond would be significantly shorter than the Ni2+ –O bond. Reduction of the Ni3+ on heating in vacuum would result in lengthening of the Ni–O bonds. However as there is an averaging of the data in the diffraction experiment the small quantity of Ni3+ present is unlikely to have a dramatic effect but will contribute to the overall bond length changes. Furthermore there is debate as to whether the Ni3+ species exists in these materials and that the charge compensation is through incorporation of O− species [22] or, depending on delta value, a combination of the Ni3+ , O− and O2− species. Hence the situation would be further complicated by bonding to the O− species. Therefore during the mild reduction of the La2 NiO4+δ
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423
Fig. 2. Neutron diffraction time-of-flight data for La2 NiO4+δ recorded at 150 ◦ C. (+) represent experimental data points, solid line is refined data fit. The upper markers are the minority NiO phase, lower markers are La2 NiO4+δ tetragonal phase. Rp = 4.13%, Rwp = 4.77% (d-spacing range 0.6–2.5 Å).
Table 2 Refined atomic positions for La2 NiO4+δ in I 4/mmm spacegroup. Data recorded at 15 ◦ C x
Name
y
z
Uiso × 100
Occupancy
La1 Ni1 O1 O2 O3
0 0 0 0 0
0 0 0.5 0 0.5
0.36088(9) 0 0 0.1741(2) 0.2339(24)
1.127(29) 1.075(33) 1.55(6) 4.04* 2.60(75)
1 1 1.000(9) 0.957(10) 0.053(4)
*
U11
U22
U33
U12
U13
U23
O2
5.71(10)
5.71(10)
0.71(11)
0
0
0
Table 3 Refined atomic positions for La2 NiO4+δ in I 4/mmm spacegroup at elevated temperature (600 ◦ C) La1 Ni1 O1 O2
z
Uiso × 100
Occupancy
0 0 0.5 0
0.36244(9) 0 0 0.17628(18)
1.817(30) 1.978(36) 2.568(46) 5.01*
1 1 1 0.995(10)
x
Name 0 0 0 0
y
*
U11
U22
U33
U12
U13
U23
O2
6.57(11)
6.57(11)
1.89(12)
0
0
0
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(a)
(b)
Fig. 3. Variation of (a) a cell constant and (b) c cell constant with temperature for La2 NiO4+δ . N.B.: Standard deviations lie within the data markers.
(a)
(b)
Fig. 4. Variation of (a) Ni–O1 and (b) Ni–O2 bond lengths with temperature. N.B.: Standard deviations lie within the data markers where not plotted.
(a)
(b)
Fig. 5. Variation of (a) La–O1 and (b) La–O2 bond lengths with temperature.
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425
Fig. 7. Change in lattice volume with temperature for La2 NiO4+δ . N.B.: Standard deviations lie within the data markers.
4. Conclusions
Fig. 6. Coordination environment of La in the tetragonal La2 NiO4 structure. Large spheres represent lanthanum, intermediate spheres are nickel and the small spheres are oxygen.
sample it is likely that the bond length changes observed are a summation of a variety of competing factors. Also further examination of the La–O bond distances for the lattice oxygen shows remarkable changes, particularly on cooling. Fig. 5 plots the La–O1 and La–O2 bond distances on cooling and it is evident that the structure compensates for the loss of the interstitials through lengthening and shortening of these bonds respectively, hence maintaining the lattice volume. Only selected La–O distances have been significantly influenced and it is necessary to consider the coordination environment around La in the K2 NiF4 structure to explain this, Fig. 6. It is evident from the data that on heating, the La–O1 bonds, that is those in the Ni– O plane, are relatively unchanged with values of around 2.62 Å through the entire temperature range. However on cooling the bond lengths decrease linearly with temperature. By contrast the La–O2 bonds, which could be viewed as equatorial, closely mirror the behaviour of the Ni–O2 bonds. All of the bond length changes on heating tend to have a degree of curvature to them which is a result of both the Ni–O and La–O bonds continually adapting to the loss of interstitial oxide ions. This is further supported by the completely linear behaviour of all bonds on cooling. A further remarkable feature of this system is that the volume changes very little, Fig. 7, despite the significant structural changes. Of course, there is an expansion of the volume on heating but it is observed to follow a linear course despite the departure from linearity of the lattice parameters and bond distances. It is of interest to note that the remarkable bond length changes and subsequent effect on the volume of the unit cell is not observed in the NiO phase where complete linear behaviour was followed.
Through investigation of in-situ high resolution neutron powder diffraction data for La2 NiO4+δ it has been established that only one structural transformation occurs and this is on heating. The tetragonal modification was observed in the data recorded at temperatures above 150 ◦ C and therefore it can be concluded that the transformation takes place at some point between room temperature and 150 ◦ C. The model which provides the best fit to the data at room temperature is the Fmmm orthorhombic model which then undergoes a transformation to the related I 4/mmm tetragonal structure. This structure is then maintained throughout both heating and cooling regimes. Loss of interstitial oxygen was observed to take place over the heating regime from room temperature until 450 ◦ C, at which point no interstitial oxygen position could be refined. Further studies of this material in which the high resolution diffraction data are recorded in air would be of further interest and such studies are planned to expand upon this work. Acknowledgements Thanks are due to the CCLRC for funding this work through a beamtime grant, RB11834 to enable the data to be collected at the ISIS facility. Thanks are also due to Dr Kevin Knight for his invaluable assistance in recording the high temperature in-situ data and Dr Peter Slater, University of Surrey. References [1] J.E. Millburn, M.A. Green, D.A. Neumann, M.J. Rosseinsky, J. Solid State Chem. 145 (1999) 401. [2] A. Mehta, P.J. Heaney, Phys. Rev. B 49 (1994) 563. [3] J.M. Tranquada, Y. Kong, J.E. Lorenzo, D.J. Buttrey, D.E. Rice, V. Sachan, Phys. Rev. B 50 (1994) 6340. [4] G. Burns, F.H. Dacol, D.E. Rice, D.J. Buttrey, M.K. Crawford, Phys. Rev. B 42 (1990) 10777.
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[5] J.E. Lorenzo, J.M. Tranquada, D.J. Buttrey, V. Sachan, Phys. Rev. B 51 (1995) 3176. [6] J. Jorgensen, B. Dabrowski, S. Pei, D. Richards, D. Hinks, Phys. Rev. B 40 (1989) 2187. [7] J. Rodriguez-Caravajal, M.T. Fernandez-Diaz, J.L. Martinez, J. Phys.: Condens. Matter. 3 (1991) 3215. [8] W. Paulus, A. Cousson, G. Dhalenne, J. Berthon, A. Revcolevschi, S. Hosoya, W. Treutmann, G. Heger, R.L. Toquin, Solid State Sciences 4 (2002) 565. [9] W. Paulus, A. Cousson, G. Heger, A. Revcolevschi, G. Dhalenne, S. Hosoya, Physica B 234–236 (1997) 20. [10] T. Kyomen, H. Tamura, M. Oguni, M. Itoh, K. Kitayama, J. Phys.: Condens. Matter. 9 (1997) 1841. [11] S.J. Skinner, J.A. Kilner, Solid State Ionics 135 (2000) 709. [12] V.V. Kharton, A.P. Viskup, A.V. Kovalevsky, E.N. Naumovich, F.M.B. Marques, Solid State Ionics 143 (2001) 337. [13] V.V. Kharton, A.P. Viskup, E.N. Naumovich, F.M.B. Marques,
J. Mater. Chem. 9 (1999) 2623. [14] M.P. Pechini, US Pat. No. 3,330,697 (1967). [15] S. Bhavaraju, J.F.D. Carlo, D.P. Scarfe, A.J. Jacobson, D.J. Buttrey, Solid State Ionics 86–88 (1996) 825. [16] P.J. Heaney, A. Mehta, G. Sarosi, V.E. Lamberti, A. Navrotsky, Phys. Rev. B 57 (1998) 10370. [17] M. Sayagues, M. Vallet-Regi, J. Hutchinson, J. Gonzalez-Calbet, J. Solid State Chem. 125 (1996) 133. [18] H. Tamura, A. Hayashi, Y. Ueda, Physica C 258 (1996) 61. [19] H. Tamura, A. Hayashi, Y. Ueda, Physica C 216 (1993) 83. [20] D.J. Buttrey, P. Ganguly, J.M. Honig, C.N.R. Rao, R.R. Schartman, G.N. Subbanna, J. Solid State Chem. 74 (1988) 233. [21] R.B. Von Dreele, A.C. Larson, General Structural Analysis System, Los Alamos National Laboratory, Los Alamos, 2001. [22] L. Minervini, R.W. Grimes, J.A. Kilner, K.E. Sickafus, J. Mater. Chem. 10 (2000) 2349.
Characterisation of La2NiO4+δ using in-situ high temperature neutron powder diffraction Stephen J. Skinner Centre for Ion Conducting Membranes, Department of Materials, Imperial College London, Prince Consort Road, London, SW7 2BP, UK Received 14 October 2002; received in revised form 16 December 2002; accepted 8 January 2003
Abstract La2 NiO4+δ , has been studied using in-situ high temperature neutron diffraction over a temperature range of 25–800 ◦ C in vacuum. The behaviour of this material, and in particular the oxygen interstitial content, is discussed and quite remarkable bond length changes observed. It is observed that at temperatures above 150 ◦ C La2 NiO4+δ transforms to the tetragonal I 4/mmm structure and maintains this over the entire temperature range on both heating and cooling. The loss of the interstitial oxygen was observed over the low temperature region of the study and significant changes in both lattice constant and bond lengths found to mirror these changes, indicating the structural importance of the interstitial oxygen. 2003 Éditions scientifiques et médicales Elsevier SAS. All rights reserved. Keywords: In-situ; Neutron diffraction; High temperature; Interstitial; K2 NiF4 ; La2 NiO4+δ
1. Introduction Although the structure of La2 NiO4+δ has been extensively investigated [1–7] there have been no in-situ high temperature studies performed, either by XRD or neutron diffraction. It has been reported that materials of the K2 NiF4 structure type can accommodate a wide variety of oxygen stoichiometries, predominantly hyperstoichiometry. This has been found to influence the symmetry of La2 NiO4+δ , with a variety of structural modifications based on the tetragonal I 4/mmm spacegroup reported. La2 NiO4+δ adopts the K2 NiF4 structure but with a slight orthorhombic distortion of the unit cell in the ab plane [6]. The stoichiometric material can be easily oxidised with the excess oxygen occupying interstitial sites and, depending on the amount of oxygen incorporated, the structure will vary between related tetragonal and orthorhombic structures. Several studies, including some utilising neutron diffraction techniques [5–8], have accurately determined the degree of orthorhombic distortion and the position of the oxygen interstitials. However there remains considerable debate as to the exact nature of the low temperature form of La2 NiO4 with both orthorhombic and tetragonal unit cells having been proposed, as well as biphasic mixtures. Indeed it has been proE-mail address: [email protected] (S.J. Skinner).
posed that a two phase region of the phase diagram exists for compositions with 0.02 δ 0.13 [6] and further details of oxygen intercalated samples indicate three two phase regions existing between 0.03 δ 0.13 [3]. The Fmmm phase has been proposed as being stable above the δ = 0.13 limit. Synthesis method has also been suggested as an influencing factor in determining the phase produced [7]. Unfortunately to date there has been little investigation of the behaviour of these oxides in-situ at elevated temperatures. This may partly be due to the lack of interest in high temperature applications of K2 NiF4 type oxides as the majority of efforts in recent years have been directed towards the La2 CuO4 related superconducting oxides and hence most studies have concentrated on the materials at < 25 ◦ C. One study, of La2 NiO4.13 , has been reported [9] utilising neutron diffraction techniques but the details of the report are limited to a maximum temperature of 400 ◦ C. A subsequent study by Paulus et al. [8] details structural refinements on single crystal samples at three temperatures, below 25 ◦ C, at 25 ◦ C and at 400 ◦ C where the materials have been electrochemically oxidised. A further study [10] has endeavoured to use heat capacity anomalies in a single crystal to identify structural changes, but as previously the study was limited, with a maximum temperature of 227 ◦ C. However, there has been increasing interest in K2 NiF4 type oxides as possible oxide ion conducting ceramics since the discovery
1293-2558/03/$ – see front matter 2003 Éditions scientifiques et médicales Elsevier SAS. All rights reserved. doi:10.1016/S1293-2558(03)00050-5
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S.J. Skinner / Solid State Sciences 5 (2003) 419–426
of mixed ionic-electronic conduction in the La2 NiO4+δ material [11–13]. Ion conducting ceramics are widely used in high operating temperature devices such as solid oxide fuel cells and ceramic oxygen generators. It is therefore essential to understand the stability and properties of these oxides at elevated temperatures if La2 NiO4+δ and related materials are to be considered as new mixed ionic-electronic conducting components in electrochemical devices such as solid oxide fuel cells. In this paper the phase stability and transitions of La2 NiO4+δ is examined and the behaviour on both heating and cooling in vacuum investigated.
2. Experimental La2 NiO4+δ was prepared using a modified Pechini [14] method in which solutions of the appropriate nitrates were combined with excess citric acid. This mixture was then heated until gelation occurred and then dehydrated and decomposed at 500 ◦ C. After decomposition the material was sintered at 1370 ◦ C for 12 h after which a single phase, as determined by XRD, was identified. In order to establish the oxygen stoichiometry of the phase produced thermogravimetric analysis in a 10% H2 /N2 mixture was performed using a Stanton Redcroft 780 series instrument. The reduction of the oxygen excess material to the component oxides and metals allowed the accurate determination of the value of δ. The in-situ neutron diffraction measurements were carried out on the HRPD beamline at ISIS at the Rutherford Appleton Laboratories. Measurements were carried out in the vacuum conditions of the apparatus to enable a monitoring of the loss of excess oxygen as temperature increased. Further, to enable any structural modifications resulting from oxygen loss to be detected both heating and cooling cycles were recorded. Heating of the samples was carried out using an induction furnace which would easily achieve the 800 ◦ C upper temperature required.
3. Results and discussion Considerable debate surrounds the crystal structure adopted by La2 NiO4+δ both at room temperature and at elevated temperatures [2,3,5–8,15–20] and in order to elucidate the structure at elevated temperature the high resolution powder diffraction (HRPD) instrument was used to investigate these phases. Data was recorded at temperatures of up to 800 ◦ C, a temperature which was deemed to be sufficient to allow the evolution of the interstitial oxygen without causing the reduction of the material to the component oxides and metals. To determine the oxygen stoichiometry thermogravimetric analysis was carried out under the previously described reducing conditions and an oxygen excess content of
0.18 ± 0.01 was determined for the as synthesised material at room temperature. Initial neutron diffraction data recorded at room temperature and in vacuum indicated that a small amount of NiO was present in the sample (∼ 2%) which was not detected in the powder XRD patterns. Due to the relatively strong scattering of the Ni ion the small amount of NiO detected is quite prominent in the neutron powder diffraction pattern. The peaks observed that were attributable to NiO were clearly distinguishable from those due to the La2 NiO4+δ phase. Obviously this observation would have an effect on the TGA results, but as the NiO would not reduce until the same temperature as the La2 NiO4 the initial weight loss at ∼ 400 ◦ C would be attributable to the interstitial loss. Therefore the small percentage of NiO present in the sample has a minimal effect on the determination of the oxygen nonstoichiometry. Subsequent structural refinements using the GSAS software [21] were therefore two component models including the Fm3m cubic NiO phase with lattice parameters of a ∼ 4.18 Å. In the initial refinements the La:Ni ratio was refined but found to be as the starting stoichiometry suggested and therefore in later refinements the occupancy of these sites was fixed at 1. The presence of the NiO phase did have a benefit to the refinement process in that the lattice parameter expansion of the oxide phase provided a useful guide in terms of the thermal expansion effect on the lattice of both materials. A completely linear behaviour on both heating and cooling was observed which indicated that the data collection processes were reliable and any changes observed in the La2 NiO4+δ diffraction patterns would be due to changes within the material and not instrumental. Several authors have proposed models for the structure of La2 NiO4+δ at room temperature and include both single phase and two phase models depending on the oxygen stoichiometry. Therefore Rietveld refinements of several possible models were initially attempted using the GSAS suite of programmes [21] and including both tetragonal and orthorhombic possibilities. A shifted Cherbyshev background function and pseudo Voigt peak shape were used in the refinement process. Early in this process two phase models were discounted as there was no evidence to support this from the experimental data through either peak shape, broadening or splitting. Contrary to the findings of Paulus [9] and Mehta [2], models fitting to space groups Bbcm and F 4/mmm were found to give unrealistic fits to the experimental data. Furthermore, attempts to refine the data to the tetragonal I 4/mmm space group resulted in a physically implausible model. A best fit to the room temperature data was obtained in space group Fmmm with the La occupying the (00z) position, Ni occupying the (000) position and oxygen occupying four positions with variable fractional occupancies, as detailed in Fig. 1 and Table 1. The refined data also enabled the oxygen content to be determined with some accuracy and it was found to be in good agreement with the TGA data, giving a final refined value over all oxygen positions of δ = 0.17. The
S.J. Skinner / Solid State Sciences 5 (2003) 419–426
421
Fig. 1. Neutron diffraction time-of-flight data for La2 NiO4+δ recorded at 25 ◦ C. (+) represent experimental data points, solid line is refined data fit. The upper markers are the minority NiO phase, lower markers are La2 NiO4+δ orthorhombic phase. Rp = 3.92%, Rwp = 4.57% (d-spacing range 0.6–2.5 Å). Table 1 Refined atomic positions for La2 NiO4+δ in Fmmm spacegroup. Data recorded at 25 ◦ C Name La1 Ni1 O1 O2 O3 O4
x
y
z
0 0 0.25 0.25 0 −0.0661(17)
0 0 0.25 0.25 0 −0.0627(16)
0.36113(7) 0 0 0.22674(167) 0.17421(27) 0.16990(50)
Occupancy 1 1 1.0031 0.038(4) 0.531(20) 0.119(6)
occupancy of the 4 oxygen positions is also in reasonable agreement with the work of Jorgensen et al. [6] who refined the room temperature structure of La2 NiO4+δ in the Fmmm spacegroup. The standard deviation for the oxygen stoichiometry is subject to a reasonably high degree of uncertainty due to the fitting of the O2 position in particular as the error in the site occupancy is of the order of 10% which is related to the low occupancy of this site, 0.038. It is also interesting to note that the O2 position has significant variation in the positional parameters as well as the thermal parameters. This uncertainty is also likely to be due to the virtually zero value for the site occupancy. This contrasts
U11 0.81(2) 0.75(2) 1.31(14) 0.2 0.21(14) 0.67(22)
U22 0.81(2) 0.75(2) 0.50(13) 0.2 0.21(14) 0.67(22)
U33 0.81(2) 0.75(2) 2.69(9) 0.2 0.21(14) 0.67(22)
U12 0 0 −0.06(5) 0 0 0
with the accuracy of the determination of the position of the fully occupied site, O1, which has a standard deviation of 0.08%. Hence, it is indicative of the oxygen stoichiometry of the sample and when used in conjunction with the data obtained from the TGA measurements can confirm the total stoichiometry of the sample and the overall oxygen losses from the material. Consequently the structure of this sample at room temperature is found to be orthorhombic Fmmm with refined lattice parameters of a = 5.45869(4) Å, b = 5.46360(4) Å and c = 12.68524(4) Å. With the room temperature data refined and in good agreement with the literature the main focus of the work
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was on the elevated temperature data sets. Two sets of lower temperature data were collected, at 150 and 300 ◦ C, with the data then collected every 50 ◦ C from 400 until 700 ◦ C. A final set of data was recorded at 800 ◦ C and then every 100 ◦ C on cooling until 300 ◦ C. Initial analysis of the lattice parameter data on heating indicated that there was non-linear behaviour in the c parameter and very little change in the a and b parameters. The interpretation of this data was assumed to be due to the oxygen loss from the La–O layers. The temperature at which this transition was evident was between 300 and 700 ◦ C and this temperature range corresponded well to the data collected from thermogravimetric analysis. Rietveld refinements of the experimental data at each temperature were performed using GSAS and each data set carefully examined for evidence of tetragonal/orthorhombic distortions. From the data recorded at 150 ◦ C, Fig. 2, there was no evidence of orthorhombic splitting and/or line broadening indicative of an orthorhombic distortion. The model used in this refinement was based on the I 4/mmm space group with an oxygen excess position at the (0 0.5 0.23) position. The refinement details are given in Table 2. In this model the anisotropic temperature factors were refined only for O2 as the refinement of this position isotropically did not give a physically plausible model. Refinement of all other temperature factors using an isotropic model, as observed in Table 2, resulted in temperature factors with reasonable values. The values for O2 gave a particularly high Uiso and on subsequent analysis of the anisotropic temperature factors it is evident that this was due to the large difference between the U11 /U22 values and that for U33 . The large ab plane anisotropy of the apical oxygen position in K2 NiF4 type oxides is well known and indeed may contribute to the fast oxide ion conduction in these materials, and is illustrated by the fitting of the temperature factors of the O2 position. To ensure that the spacegroup assignment was correct data refinements were performed in the orthorhombic spacegroups detailed previously but there were no significant improvements in the fits or indeed in the Rwp and Rp values. From the refinement in the tetragonal spacegroup it was possible to determine an oxygen excess stoichiometry of δ ≈ 0.13. Further refinements up to 800 ◦ C revealed no further structural transformations and all data could be refined in the I 4/mmm space group. Typical data from a refinement at 600 ◦ C are given in Table 3. Only the change in oxygen stoichiometry was evident when the lattice parameters were examined and the oxygen positions refined. As anticipated from previous studies the overall oxygen excess concentration determined from the Rietveld refinement of the oxygen positions decreased from ∼ 4.17 at room temperature to ∼ 4.00 at 450 ◦ C. Examination of the variation in the lattice parameter data with temperature, Fig. 3, is revealing. On heating there is in fact very little change in both the a and b parameters (orthorhombic lattice parameters from the room temperature
√ data have been converted through a 2a relationship to the tetragonal cell) but on studying the behaviour of the c parameter it is obvious that there is a large change associated with the loss of oxygen. It is also telling that at both the low and high temperature regions the c parameter behaviour is linear, whereas there are significant differences in the behaviour around the transition temperature range. As the oxygen stoichiometry has undergone significant change at the lower temperature it is interesting to investigate the reason for the linear behaviour of the c parameter in this region. In the transition region there are two main competing demands on the structure—adjustment to the loss of interstitial oxygen and thermal expansion. The linearity of the expansion in the c cell constant at temperatures below 400 ◦ C and above 650 ◦ C can be attributed to thermal expansion whereas between 400 and 650 ◦ C there is the additional loss of oxygen modifying the structure. These competing demands explain the deviation from linearity in the lattice parameters with temperature and also explain the linearity of the plot on cooling, where there is no variation in oxygen stoichiometry. It is also well known [2,6] that the oxygen interstitial site lies in the La–O plane in the ab direction close to the (0.25, 0.25, 0.25) position and therefore further evidence for the loss of oxygen from interstitial sites should be evident from the bond distance data. Furthermore, these reports indicate that the incorporation of the interstitial oxygen forces the apical oxygen to be forced into two split positions (0, 0, ∼ 0.17) and (−0.06, −0.06, 0.17) and the resultant effect on bond distance should be evident. Looking at the detail of the bond distances it is of interest to note the behaviour of the equatorial and apical Ni–O bonds in the NiO6 octahedron. It was anticipated that the Ni–O bonds would lengthen with increase in temperature and indeed this was observed for both the equatorial and apical positions, Fig. 4. However, on closer examination it is evident that the two apical Ni–O bonds have experienced a much greater effect than the equatorial positions. It is also interesting to observe that the apical position bond lengths undergo a dramatic lengthening at ∼ 400–450 ◦ C, indicating that this change was directly attributable to the loss of interstitial oxygen ions. It is likely that in the oxygen excess materials there is a small quantity of Ni3+ present. If this is the case then the Ni3+ –O bond would be significantly shorter than the Ni2+ –O bond. Reduction of the Ni3+ on heating in vacuum would result in lengthening of the Ni–O bonds. However as there is an averaging of the data in the diffraction experiment the small quantity of Ni3+ present is unlikely to have a dramatic effect but will contribute to the overall bond length changes. Furthermore there is debate as to whether the Ni3+ species exists in these materials and that the charge compensation is through incorporation of O− species [22] or, depending on delta value, a combination of the Ni3+ , O− and O2− species. Hence the situation would be further complicated by bonding to the O− species. Therefore during the mild reduction of the La2 NiO4+δ
S.J. Skinner / Solid State Sciences 5 (2003) 419–426
423
Fig. 2. Neutron diffraction time-of-flight data for La2 NiO4+δ recorded at 150 ◦ C. (+) represent experimental data points, solid line is refined data fit. The upper markers are the minority NiO phase, lower markers are La2 NiO4+δ tetragonal phase. Rp = 4.13%, Rwp = 4.77% (d-spacing range 0.6–2.5 Å).
Table 2 Refined atomic positions for La2 NiO4+δ in I 4/mmm spacegroup. Data recorded at 15 ◦ C x
Name
y
z
Uiso × 100
Occupancy
La1 Ni1 O1 O2 O3
0 0 0 0 0
0 0 0.5 0 0.5
0.36088(9) 0 0 0.1741(2) 0.2339(24)
1.127(29) 1.075(33) 1.55(6) 4.04* 2.60(75)
1 1 1.000(9) 0.957(10) 0.053(4)
*
U11
U22
U33
U12
U13
U23
O2
5.71(10)
5.71(10)
0.71(11)
0
0
0
Table 3 Refined atomic positions for La2 NiO4+δ in I 4/mmm spacegroup at elevated temperature (600 ◦ C) La1 Ni1 O1 O2
z
Uiso × 100
Occupancy
0 0 0.5 0
0.36244(9) 0 0 0.17628(18)
1.817(30) 1.978(36) 2.568(46) 5.01*
1 1 1 0.995(10)
x
Name 0 0 0 0
y
*
U11
U22
U33
U12
U13
U23
O2
6.57(11)
6.57(11)
1.89(12)
0
0
0
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(a)
(b)
Fig. 3. Variation of (a) a cell constant and (b) c cell constant with temperature for La2 NiO4+δ . N.B.: Standard deviations lie within the data markers.
(a)
(b)
Fig. 4. Variation of (a) Ni–O1 and (b) Ni–O2 bond lengths with temperature. N.B.: Standard deviations lie within the data markers where not plotted.
(a)
(b)
Fig. 5. Variation of (a) La–O1 and (b) La–O2 bond lengths with temperature.
S.J. Skinner / Solid State Sciences 5 (2003) 419–426
425
Fig. 7. Change in lattice volume with temperature for La2 NiO4+δ . N.B.: Standard deviations lie within the data markers.
4. Conclusions
Fig. 6. Coordination environment of La in the tetragonal La2 NiO4 structure. Large spheres represent lanthanum, intermediate spheres are nickel and the small spheres are oxygen.
sample it is likely that the bond length changes observed are a summation of a variety of competing factors. Also further examination of the La–O bond distances for the lattice oxygen shows remarkable changes, particularly on cooling. Fig. 5 plots the La–O1 and La–O2 bond distances on cooling and it is evident that the structure compensates for the loss of the interstitials through lengthening and shortening of these bonds respectively, hence maintaining the lattice volume. Only selected La–O distances have been significantly influenced and it is necessary to consider the coordination environment around La in the K2 NiF4 structure to explain this, Fig. 6. It is evident from the data that on heating, the La–O1 bonds, that is those in the Ni– O plane, are relatively unchanged with values of around 2.62 Å through the entire temperature range. However on cooling the bond lengths decrease linearly with temperature. By contrast the La–O2 bonds, which could be viewed as equatorial, closely mirror the behaviour of the Ni–O2 bonds. All of the bond length changes on heating tend to have a degree of curvature to them which is a result of both the Ni–O and La–O bonds continually adapting to the loss of interstitial oxide ions. This is further supported by the completely linear behaviour of all bonds on cooling. A further remarkable feature of this system is that the volume changes very little, Fig. 7, despite the significant structural changes. Of course, there is an expansion of the volume on heating but it is observed to follow a linear course despite the departure from linearity of the lattice parameters and bond distances. It is of interest to note that the remarkable bond length changes and subsequent effect on the volume of the unit cell is not observed in the NiO phase where complete linear behaviour was followed.
Through investigation of in-situ high resolution neutron powder diffraction data for La2 NiO4+δ it has been established that only one structural transformation occurs and this is on heating. The tetragonal modification was observed in the data recorded at temperatures above 150 ◦ C and therefore it can be concluded that the transformation takes place at some point between room temperature and 150 ◦ C. The model which provides the best fit to the data at room temperature is the Fmmm orthorhombic model which then undergoes a transformation to the related I 4/mmm tetragonal structure. This structure is then maintained throughout both heating and cooling regimes. Loss of interstitial oxygen was observed to take place over the heating regime from room temperature until 450 ◦ C, at which point no interstitial oxygen position could be refined. Further studies of this material in which the high resolution diffraction data are recorded in air would be of further interest and such studies are planned to expand upon this work. Acknowledgements Thanks are due to the CCLRC for funding this work through a beamtime grant, RB11834 to enable the data to be collected at the ISIS facility. Thanks are also due to Dr Kevin Knight for his invaluable assistance in recording the high temperature in-situ data and Dr Peter Slater, University of Surrey. References [1] J.E. Millburn, M.A. Green, D.A. Neumann, M.J. Rosseinsky, J. Solid State Chem. 145 (1999) 401. [2] A. Mehta, P.J. Heaney, Phys. Rev. B 49 (1994) 563. [3] J.M. Tranquada, Y. Kong, J.E. Lorenzo, D.J. Buttrey, D.E. Rice, V. Sachan, Phys. Rev. B 50 (1994) 6340. [4] G. Burns, F.H. Dacol, D.E. Rice, D.J. Buttrey, M.K. Crawford, Phys. Rev. B 42 (1990) 10777.
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