Controlling parameters for superconductivity in layered lithium iron hydroxide selenides.† Hualei Sun,a,b Daniel N. Woodruff,a Simon J. Cassidy,a,c Genevieve M. Allcroft,a Stefan J. Sedlmaier,a Amber L. Thompson,a Paul A. Bingham,d Susan D. Forderd Simon Cartenet,d Nicolas Mary,d Silvia Ramos,e Francesca R. Foronda,f Benjamin H. Williams,f Xiaodong Li,b Stephen J. Blundellf and Simon J. Clarke*a Hydrothermal synthesis of layered lithium iron selenide hydroxides Li1–xFex(OH)Fe1–ySe (x ~ 0.2; 0.03 < y < 0.15) is used to open up further soft chemical routes to iron-based superconductors. Control of the hydrothermal reactions and post-synthetic compositional control by lithiation to avoid iron deficiency in the selenide layers shows that the superconducting properties in these compounds are enhanced by minimising the concentration of iron vacancies in the iron selenide layer and simultaneously increasing the electron count on iron in the selenide layers.
Introduction Iron-based superconductors1 form a class of unconventional superconductor in which superconductivity at temperatures as high as 50 K emerges from close proximity to a competing itinerant antiferromagnetic state. Control over the physical properties, particularly the competition between the antiferromagnetic state and the superconducting state, may be exerted by making isovalent or aliovalent substitutions,1,2 often over very narrow compositional ranges,3,4 or by applying an external pressure.5,6 Bond lengths and angles7 in the iron arsenide, phosphide or selenide layers and the electron count1,2,3,4 (formal iron oxidation states are around Fe2+) are established as key controlling parameters. Unravelling the compounds’ behaviour may shed light on the behaviour of other classes of high temperature superconductor, such as the layered cuprates, which currently defy a full theoretical understanding. Recent reviews describe developments in synthesis,8,9 structure-property relationships,10 materials properties11 and theory.12 The tetragonal polymorph of iron selenide is a superconductor with a superconducting critical temperature, Tc, of 8.5 K when close to stoichiometric (~1 % additional Fe in interstitial sites between the FeSe layers is intrinsic).13,14 Derivatives of this compound and thin films15 have recently been developed as higher temperature superconductors with Tcs exceeding 30 K. These include high-temperature-synthesised phases (~900 °C) A1– 16 xFe2–2ySe2 (x ~ y ~ 0.2; the so-called 2-4-5 phases) with a highly defective ThCr2Si2-type structure. In these samples the bulk of the sample is a Mott-Hubbard insulator with a large local moment on Fe and an ordered array of iron site vacancies,17,18 but superconductivity with Tc of about 30 K is often found in minority regions of the sample with less iron deficiency and less alkali metal.19,20,21 In order to decrease the concentration of iron site vacancies in the FeSe layers, stoichiometric, superconducting
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FeSe itself has been used in the synthesis, at ambient temperatures and below, of intercalates using solutions of electropositive metals in ammonia.22 These intercalates, which often superconduct at temperatures as high as 45 K contain variable electropositive metal and ammonia and amide contents and are the subject of current investigation.23,24,25,26
Figure 1. The structure of the lithium iron selenide hydroxides and a typical refinement against neutron diffraction data from the HRPD instrument at the ISIS facility showing data (blue dots), calculated (red line), difference (black line) and reflection positions. Data for the 168° bank have been displaced by 9 units along the vertical axis. Further refinements are supplied in Figures S1 and S2.
In this article we describe the use of tetragonal FeSe as a precursor in hydrothermal synthesis at modest temperatures (about 200 °C) and highly basic conditions (to avoid hydrolysis to H2Se) of layered hydroxide selenides (Figure 1). We identify the iron content in the selenide layer, and the related iron oxidation state as the key variables controlling the details of the crystal structure, whether the samples superconduct, and the value of the superconducting Tc. The highest Tc compounds in this series (Tc ~ 41 K) are obtained by reducing the concentration of vacancies on the iron site in the selenide layers using a post-synthetic lithiation treatment. A similar hydrothermal route using different precursors and a subsequent high-pressure annealing step was reported by Lu et al. to produce a related superconducting oxide selenide.27 But this compound has subsequently been shown by the same authors28 to be a member of the series of hydroxide selenides reported here.
Synthesis and preliminary analysis Hydrothermal synthesis Hydrothermal syntheses were carried out in stainless steel autoclaves with teflon liners of about 18 cm3 capacity in a procedure modified from that described in reference 27. The Pourbaix diagram for iron and selenium is known from investigations of the contamination of natural waters29 and reveals that under reducing conditions and at high pH values the formation of H2Se is suppressed and FeSe is stable. This accords with routes to other iron hydroxide chalcogenides such as tochilinite30 in which iron sulfide layers are separated by magnesium hydroxide layers. In a typical synthesis 6 mmol (0.8 g of tetragonal FeSe (synthesised from the elements (Fe ALFA 99.998 %; Se ALFA 99.999%) as described previously14), 140 mmol (6 g) of LiOH·H2O (Aldrich 98%) and 5 ml of deionised and de-oxygenated water were loaded into the autoclave together with variable amounts of additional iron powder. The autoclaves were tightly sealed and placed in a chamber furnace. The furnace was heated to 200 °C at 1°C per minute and the temperature was maintained for 12 days. The furnace was then turned off and allowed to cool naturally and the autoclaves were removed at room temperature. The autoclaves were opened in an argon-filled glove bag and the products were loaded into Schlenk tubes prior to washing three times with deionised and de-oxygenated water to remove soluble side products. Magnetic impurities were removed from some syntheses using a strong magnet. The samples were then dried under vacuum and removed to an argonfilled glovebox. The highly crystalline samples were black with metallic lustre and were examined with no further synthetic treatment. As described below, superconducting samples were always obtained when the overall ratio of Fe:Se in the synthesis was 1.16:1 (i.e. 1 mmol additional Fe for 6 mmol FeSe in the autoclave). The use of smaller amounts of additional Fe (0 or 0.5 mmol Fe per 6 mmol of FeSe) produced non-superconducting products which were of high quality as judged by powder diffraction measurements. The use of larger amounts of additional Fe led to significant contamination by iron oxide side products. The synthesis was found to be scalable to produce 10 g of product by increasing the amount of FeSe and Fe in the synthesis 12 fold, increasing the amount of water to 7 cm3 and
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maintaining the amount of LiOH which remains in a large excess. Some of these samples were subsequently subjected to lithiation in which the powders were stirred in solutions of lithium in liquid ammonia at –30°C using a Schlenk line, with subsequent evaporation of the solvent and evacuation to yield the dried product. Magnetometry Measurement of all the samples using SQUID magnetometry (Quantum Design MPMS-XL magnetometer) using measuring fields of 20 – 50 Oe, and with the samples sequestered from air revealed that several of the samples showed superconductivity with Tcs in the range 10 – 39 K and with variable shielding fractions. However the samples made with the smaller Fe:Se ratios in the synthesis mixture did not exhibit superconductivity. Figure 2(a) shows the variation in the magnetic properties for several samples representing the full range of properties in the hydrothermally-synthesised samples.
Figure 2. (a) Superconductivity in the hydrothermally synthesised samples showing the range of superconducting Tcs. Zero-field-cooled (filled symbols) and field-cooled (open symbols) data are shown. Some samples showed a high normal state background due to minuscule amounts of magnetic impurities. The sample with the highest Tc exhibits diamagnetism in the field-cooled measurement, but also shows a low temperature transition below 20 K which is presumed to be due to a magnetic impurity present at levels below the detection limit of our diffraction experiments. The SI conventions were used in determining the dimensionless magnetic susceptibility which was corrected for the effect of demagnetising fields arising from the shape of the sample.31 (b) Plot of the 001 reflection measured for a range of samples on I11 showing the correlation between the c lattice parameter and whether the compounds are superconducting.
X-ray powder diffraction Laboratory X-ray powder diffraction (XRPD) investigations (Panalytical X’Pert Pro diffractometer using CuK radiation) revealed that phase pure materials were synthesised in the reactions carried out over 12 days with subsequent washing, while shorter reaction times led to the presence of oxide impurity phases. The powder diffractograms resembled that reported in ref 27 and were readily indexed on a tetragonal cell in P4/nmm with approximate lattice parameters of a = 3.8 Å and c = 9.2 Å. Higher resolution synchrotron XRPD carried out on beamline I1132 at the Diamond Light Source, UK using 0.8 Å X-rays and the multi analyser crystal detector bank revealed that the lattice parameters of the range of samples synthesised from different compositions covered a range of values (Tables S1 & S2). The values of the basal lattice parameter a spanned a range of 0.9 % and the values of the lattice parameter c spanned a range of 1.1 % (variation in the position of the 001 reflection is shown in Figure 2(b)). Small values of a corresponded to large values of c, so the unit cell volumes spanned just 0.8 %. Samples with unit cell volumes < 133.2 Å3 and c/a >2.43 (i.e. long c axes and short a axes) obtained from the iron-rich syntheses were superconducting while those with cell volumes of >133.2 Å3 and c/a < 2.43 obtained from the iron-poor syntheses were non-superconducting. Air sensitivity The samples were synthesised under reducing hydrothermal conditions obtained by incorporating elemental iron into the syntheses along with FeSe, and by excluding oxygen from the synthesis. The washing and storage of the samples was also conducted under anaerobic conditions. Exposure of the samples to air for one week was found to broaden the superconducting transition, and reduce the shielding fraction, as shown in Figure S4. However the superconducting state was not completely destroyed. This is discussed further below.
Correlation of composition, structure and properties for hydrothermally synthesised samples X-ray single crystal and neutron powder diffraction Close inspection of the hydrothermally synthesised samples revealed the presence of readily-cleavable square plate-like crystals with approximate dimensions 10 × 10 × 1 m (see Figure S5) suitable for single crystal X-ray diffraction (SCXRD) investigations using a high flux source, and these were performed at 100 K on the beamline I19 at Diamond33 using 0.68890 Å Xrays. The results of these SCXRD measurements were used with the synchrotron powder X-ray diffraction measurements and neutron powder diffraction (NPD) investigations performed on the GEM and HRPD Diffractometers at the ISIS facility, UK to explore the structural variation implied by the range of observed lattice parameters and correlate the compositions and structures with the variation in superconducting properties. The refinements against SCXRD were conducted using the CRYSTALS software,34 and the refinements against powder diffraction data were conducted using TOPAS Academic.35 The initial structural model for the samples synthesised by the hydrothermal method was obtained initially from PXRD data
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with reference to ref 27 and confirmed by solving the structure from the SCXRD data ab initio using SuperFlip36 implemented within CRYSTALS. Parallel analysis of the NPD data (see below) enabled the deduction of a final structural model in space group P4/nmm which was then refined against each single crystal dataset in turn before weights were optimised and the structure refined again. In each case, 6 cycles led to convergence. We discuss the features of the structure with reference to the diagram in Figure 1. Iron selenide layers similar to those in FeSe itself (i.e. with the anti-PbO structure type) are separated by “spacer” layers which have a similar topology, but are flattened along the c direction. Using SCXRD, all the atoms (two per selenide in total and occupying the two distinct sites labelled O and Li/Fe in Figure 1) in this “spacer” layer could be approximated by oxygen atoms with full occupancy. The shortest distances between the selenide ions in the FeSe layer and the atoms labelled O in the “spacer” layer were 3.62 Å, only marginally shorter than the interlayer Se∙∙∙Se distances of 3.71 Å in tetragonal FeSe14 and marginally longer than the between-layer Se∙∙∙Se distance of 3.58 Å in TiSe2,37 and thus longer than one would expect for Se∙∙∙O nonbonded distances. Analysis of the NPD data enabled a chemically sensible model to be obtained. These data revealed two key structural features: first that the sites labelled Li/Fe in the centres of the puckered “spacer” layers separating the FeSe layers were approximately null scattering, and second that an additional region with a negative scattering density corresponding to an approximately fully occupied hydrogen atom (coherent scattering length, b = –3.74 fm38) was located about 1 Å from the atoms in the outer part of the “spacer” layers labelled as O in Figure 1. Using data from neutron and X-ray diffraction together for multiple samples, covering the range of observed lattice parameters, produces an unambiguous structural model in which lithium iron hydroxide layers separate the iron selenide layers as shown in Figure 1. The Li:Fe ratio on the site in the centre of the hydroxide layers was consistently found to be 0.8 : 0.2 such that the site is approximately null scattering for neutrons (bLi = –1.90 fm; bFe = 9.45 fm)38 and has the average electron count similar to that of oxygen. The composition of the compounds was determined to be Li1−xFex(OH)Fe1−ySe with x ~ 0.2 and y representing a small deficiency on the Fe1 site in the iron selenide layers of about 5 – 10 % which is quantified in more detail below. Semiquantitative Energy Dispersive analysis of Xrays (EDX) conducted using an FEI Quanta 650 FEG SEM equipped with an Oxford Instruments Aztec EDS detector produced Fe:Se ratios of about 1.1 : 1 with a 3–5 % uncertainty, consistent with the composition obtained from the crystallographic measurements and poorer in iron than the 1.5 : 1 ratio proposed for hydrothermally synthesised compounds described in ref. 27, but in line with ref. 28. The presence of the H atom, forming part of a hydroxide ion, accounts for the unusually large separation of the oxide and selenide layers and the Se–H distances of about 3.1 Å correspond well to those found for weak hydrogen bonding interactions.24,39 The FeSe and Li/Fe(OH) layers are held together weakly by these hydrogen bonding interactions which accounts for the ready cleavage of the crystals. Thermogravimetric analysis carried out by heating the samples under a stream of dry N2 showed a mass loss
commencing at about 350 °C which was consistent with dehydration (see Figure S6). Detailed analysis of the SCXRD and NPD data enable the compositional and structural variation manifested in the range of lattice parameters described above to be quantified and correlated with the presence of superconductivity and with the Tc. Two samples representing the extremes of the c/a ratios found by PXRD were analysed on the GEM diffractometer at ISIS. Data were collected at room temperature and at 50 K. Comparison of the data sets gathered at the two temperatures showed no evidence for the appearance of magnetic scattering at low temperatures, nor any structural transitions. The data gathered at the two temperatures also produced similar site occupancies showing that the wide d-spacing range available on the time-offlight diffractometer, the high crystallinity, and the almost flat neutron form factor minimise parameter correlations in the refinements against these highly crystalline samples, and that this method is robust for determining site occupancy factors with a high precision. We note here that after the neutron diffraction measurements these two samples were subjected to a postsynthetic lithiation treatment and further neutron diffraction measurements were performed on the daughter products as described below. Further NPD data were collected on six additional samples at room temperature using the GEM and HRPD diffractometers and the results of these refinements are presented in Table S1. A typical refinement is shown in Figure 1 and further refinements are provided in the electronic supplementary information (Figures S2 – S3). Single crystals extracted from several of these samples were found to faithfully represent the bulk of the sample probed in the NPD experiments and the results of the refinements against I19 SCXRD data are therefore also included in the analysis that follows. The SCXRD refinements are summarised in the electronic supplementary information (Table S2).‡ The structural variable which accounts for the variation in the basal lattice parameter a is the site occupancy of the Fe1 site which is the tetrahedral site in the selenide layers (labelled Fe1 in Figure 1). Figure 3 shows the linear correlation obtained from the NPD and SCXRD measurements between the lattice parameter a and the occupancy of this site which ranges from 0.85(1) to 0.98(1) (note: the uncertainties in these parameters are those obtained from performing independent refinements against data collected at two temperatures for two of the samples; these uncertainties are about double the estimated standard deviations produced by the refinement software). Increasing the site occupancy evidently strengthens the Fe–Fe bonding within the FeSe layer, shortening the lattice parameter. There are no other significant structural differences evident in the other parts of the structure. The Li:Fe ratios on the Fe/Li site in the hydroxide layers are similar for all samples measured, and show a much weaker dependence on the a lattice parameter than the Fe1 site occupancy as Figure 3 shows. The H-contents are similar for all samples with refined occupancies between 0.95 and 1.0 for this site (note that because of the larger displacement
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ellipsoids for the H atoms determined in the NPD refinements the correlation between these and the site occupancy were slightly larger than for the other sites. We found no evidence for any hydrogen deficiency outside the uncertainty). Equivalent isotropic displacement parameters were similar for corresponding sites in the series of samples (note that in the NPD analysis this parameter was not refined for the almost null scattering Li/Fe site). The SCXRD refinements showed that the Li/Fe site had a displacement ellipsoid that was significantly elongated parallel to the c direction (U33/U11 ~ 6.5), which may reflect, evident in Figure 1, that the hydroxide layer is relatively flat compared with the selenide layer. All other ellipsoids were found to be fairly isotropic using both radiations.
Figure 3. Plot of basal lattice parameter against site occupancy of Fe in the iron selenide layer and of Li in the hydroxide layer obtained from refinements against NPD and SCXRD data. Note that the lattice parameters for the SCXRD samples were those obtained at ambient temperature using synchrotron powder XRD (beamline I11 at Diamond). The single crystal of the sample with the largest a lattice parameter had an unusually large mosaic spread which is the likely origin of the relatively large errorbars on the refined occupancies.
Key structural parameters for iron-based superconductors are the Fe−Fe distance in the plane which is equal to a / √2, the Fe−E (E = chalcogen or pnictogen) bond length and the E−Fe−E angles in the FeE4 tetrahedra. For the current compounds the shape of the FeSe4 tetrahedron is similar to that in FeSe itself and other intercalates of FeSe and is characterised by being extremely squashed in the basal plane relative to the more regular tetrahedra found in iron arsenide superconductors.7 Figure 4 shows, using normalised values of the parameters obtained from the ambient temperature NPD refinements, that the Fe−Se distance is rather invariant across the series and the effect of the change in a lattice parameter is manifested in the Se−Fe−Se angles.
a broad reflection in the neutron diffraction data at 5.565 Å (Figure S2). This was invariant in intensity with temperature suggesting that it may not be due to magnetic order, but may arise from short range structural ordering of the Li and Fe ions in the oxide layers. The non-superconducting samples with cation vacancy concentrations in the selenide layer of up to 15 % did not exhibit these broad features (Figure S2) nor was there evidence in the SCXRD or NPD data for the long-range iron/vacancy order17 found in the even more iron deficient (20% vacancies) “2-4-5” A1–xFe2–2ySe2 (x ~ y ~ 0.2) phases. This is also consistent with our investigations of related compounds containing Cu1–yS layers similar to the iron-deficient Fe1–ySe layers considered here, which show that long range vacancy order occurs at vacancy concentrations of 20 % or greater, larger than those encountered in this case.40,41
Figure 4. The variation with a (=√2×Fe−Fe) of Fe−Se bond lengths (●) and the smaller Se−Fe−Se angle (often denoted ) of multiplicity two (●) normalised against the smallest value in each series. Obtained from NPD at ambient temperatures.
Figure 5 shows that whether the samples superconduct depends upon the site occupancy of the Fe1 site. When this occupancy exceeds 95 %, the samples exhibit superconductivity, while samples with lower site occupancies are non-superconducting. Because of the linear correlation between the Fe site occupancy and the lattice parameter and tetrahedral bond angles, the compounds exhibit superconductivity when the Fe−Fe distance is short (less than 2.68 Å) and the angle is small (less than 103.4°).
To summarise the structural analysis of samples prepared directly by hydrothermal methods: there is a direct correlation between the iron content in the tetrahedral sites in the selenide layer and the basal lattice parameter. The basal lattice parameter a decreases as the number of vacancies decreases reflecting an increase in the intralayer Fe−Fe bonding as the number of vacancies decreases. It is also clear that these related changes have an important effect on whether the compounds superconduct and how high the Tc is. If the iron selenide layers have more than 5% of the sites vacant the compounds do not show superconductivity. The SCXRD and NPD analysis showed no significant variations in the composition of the other crystallographic sites from sample to sample. Therefore a high value of the iron occupancy in the selenide layer corresponds to a low oxidation state for Fe. Computing the mean iron oxidation state from the refined composition for all the hydrothermallysynthesised samples probed by NPD and SCXRD shows that for iron oxidation states greater than +2 superconductivity is not observed, while reduction of iron leads to the appearance of superconductivity and Tc increases as the formal oxidation state decreases (Figure 6).
Figure 5. The correlation of superconducting Tc with refined Fe site occupancy in the selenide layers obtained from refinements against NPD and SCXRD.
The neutron diffraction data collected at ambient and low (50 K) temperatures revealed no evidence for long range magnetic order in any of the samples. The two superconducting samples with the highest site occupancies for Fe in the selenide layer both showed
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Figure 6. Correlation of superconducting Tc with Fe oxidation state obtained from the compositions refined from NPD and SCXRD measurements.
Local probes of structure and oxidation state Mössbauer spectroscopy Preliminary ambient temperature 57Fe Mössbauer spectroscopy measurements were collected on a superconducting sample (Figure 7). The Mössbauer absorption is in the form of two paramagnetic doublets, both consistent with Fe(II). The isomer shift of the more intense doublet closely resembles that found in FeSe.42 The minor component the isomer shift is consistent with high-spin Fe2+ in the hydroxide layer.43
Figure 7. Room temperature 57Fe Mössbauer spectrum of a lithium iron hydroxide selenide sample with Tc = 26 K. The major component is ascribed to Fe1 and resembles the spectrum for FeSe while the minor doublet is consistent with Fe2+ high spin ions located in the hydroxide layers.
Magnetometry in the normal state Measurements of the magnetic susceptibility made up to room temperature in the paramagnetic region produced a Curie-Weiss type dependence with Curie constants which were much larger than could be ascribed to small paramagnetic impurities. Curie Weiss fits are shown in Figure S4. The magnitude of the Curie constant corresponded to a paramagnetic contribution from S = 2 moments carried by the Fe2+ ions (tetrahedral d6) on the Li/Fe site in the hydroxide layer. Exposure of samples to laboratory air for 1 week resulted in an increase on the Curie constant consistent with oxidation of these species to Fe3+ (tetrahedral d5). This oxidative treatment was also found to lead to the partial destruction of superconductivity. X-ray absorption spectroscopy X-ray absorption spectroscopy at the Fe-K-edge (with all spectra calibrated against an Fe foil) was used as an additional probe of the Fe oxidation state and the structure. These measurements were conducted in transmission mode on beamline B18 at Diamond44 with the samples sequestered from air and diluted with cellulose powder. The data were analysed using Athena and Artemis, part of the Demeter software package.45 Previous measurements46 of the edge position for the structurally related materials NaFeAs, KFe2As2 and NaFe1.67As2 with oxidation states of +2, +2.5 and +3 respectively showed that oxidation from +2 to
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Figure 8. (a) Fe K-edge positions of iron arsenide (red symbols) and hydrothermally-synthesised Li1–xFex(OH)Fe1–ySe superconductors (filled coloured symbols) plotted as a function of Fe oxidation state obtained from the refined compositions from diffraction data. The curved arrows show the evolution of the edge positions after lithiation (open coloured symbols). The inset shows the first derivative of the XANES absorption with the curves carrying the same colour as the points in the main figure. FeSe is included for comparison in both figures. (b) Comparison of the Fe K-edge EXAFS region for a series of hydrothermally-synthesised samples (coloured solid lines; similar colours for each sample are used in (a) and (b)) and lithiated samples (coloured dotted lines) compared with FeSe and K0.8Fe1.6Se2 (fits are provided in Figures S7 & S8). The EXAFS region is extremely sensitive to the Fe content in the selenide layer as described in the text.
+3 resulted in a shift of the edge position by 1.06 eV. Analysis of the XANES region for samples of the hydrothermally synthesised hydroxide selenides, representative of the full range of a lattice parameters probed by diffraction methods, produced edge
positions, judged by the centre peak in the first derivative of the Fe K-edge absorption, which showed a spread of values spanning 0.34 eV, suggesting a spread of oxidation states spanning approximately 0.3. When the absolute values of the edge positions were plotted against the Fe oxidation state obtained from analysis of the refined composition from diffraction measurements, with a standard sample of Fe1.01Se included for calibration of the absolute value of the edge position, a reasonably linear dependence was obtained (filled coloured symbols in Figure 8(a)) and the gradient of the line was similar to that found for the variation of edge position with Fe oxidation state in the alkali metal iron arsenide series (red symbols in Figure 8(a)).46 Note that the edge shifts significantly between arsenides and selenides showing that the absolute edge position is dependent on both local coordination environment and oxidation state. The EXAFS spectra (Figure 8(b)) were found to show a sharp sample dependence and the modelling of the EXAFS (see Figure S7 & S8) revealed that this difference was dominated by the vacancy concentration in the Fe1–ySe layers as found in the comparison between FeSe and the K0.8Fe1.6Se2 phases47,48 and in related iron arsenides.46 The comparison of the EXAFS spectra at around R = 2–3 Å for FeSe and the most Fe-poor Li1– xFex(OH)Fe1–ySe sample is similar to the comparison between FeSe and K0.8Fe1.6Se2 (Figure 8(b) shows our own sample of K0.8Fe1.6Se2, but see also Figure 2 in reference 47). Figure 8(b) shows the evolution of this region with the concentration of vacancies in the Fe1–ySe layer for the hydrothermally synthesised samples (solid lines in Figure 8(b)). Note that because Fe is present in two sites the EXAFS data at the Fe K-edge is a convolution of the contributions of both sites. Refinement against the Fe K-edge EXAFS data (Figures S7 & S8) produced ratios of iron on the Fe1 site in the selenide layer and on the Li/Fe site in the hydroxide layer that were consistent with the values obtained by analysis of the diffraction data (Table S4).
Turning on superconductivity via post synthetic lithiation. The structural analysis of the as-made hydrothermally synthesised samples shows that for superconductivity to be observed in these compounds the occupancy of the iron site should be high which results in a low oxidation state for iron. A subsequent reductive synthetic step was then applied to two of the as-made samples (one non-superconducting and the other superconducting) which had been investigated by neutron diffraction on the GEM instrument. The powders were reacted with solutions containing 0.5 moles of Li (an excess) dissolved in ammonia per mole of selenide ions in the hydroxide selenides. The results of these reactions were that the main phase seemed to retain its structure, but with a significant increase in the interlayer cell parameter c and a concomitant decrease in the basal lattice parameter a (summarised in Table 1). When the sample of the non-superconducting hydroxide selenide with 13(1) % vacancies in the Fe1–ySe layers was subjected to this treatment the c lattice parameter increased by 4.3% and the basal a lattice parameter decreased by 1.6 %, but when the superconducting sample with a 4(1)% deficiency on the Fe1 site was so-reacted, the increase in
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the c parameter (0.97 %) and the decrease in the a parameter (0.35 %) were both much smaller. In both cases the products of the post synthetic treatment were found to be superconductors with large volume fractions and Tcs of 41 K and 42 K respectively (Figure 9), larger than in any of the as-synthesised hydrothermal samples. The lithiated samples were found to be air sensitive with only vestigial superconductivity evident after 1 week of air exposure (Figure S4).
Figure 9. Changes in magnetic susceptibilities on lithiation of (a) superconducting and (b) non-superconducting hydrothermally synthesised materials. The “parent” and “daughter” designations are used in Table 1. Daughter A is depicted in Figure 10 and was used for the SR measurements (Figure 12). (c) and (d) show the corresponding shifts in the edge positions measured by XANES effected by the lithiation using the same colour scheme as in (a) and (b) respectively.
NPD data (Figure 10 and Figure S3) were collected using HRPD on these daughter samples obtained from the Li/NH3 treatment. In both cases the refined site occupancy of the Fe1 tetrahedral site in the selenide layers increased to 1.00(1). The only constraint applied to the refinements was that the Li/Fe site in the hydroxide layer was fully occupied. Comparison of both sets of parent and daughter samples (Table 1) showed that in both cases this Li/Fe shared site had an iron content which decreased by an amount equal to the increase in the occupancy of the Fe1 site (and it thus had an equally increased Li content). The refinement process did not impose a constraint on the overall Fe contents, but within both parent/daughter pair the total Fe content remained invariant after lithiation within an uncertainty in the composition similar to that found when probing a single compound at two temperatures (see above). The refined H occupancy did not change measurably under this treatment. For the superconducting sample in which the Fe1 site occupancy changes by just 4 % and the change in overall composition is fairly small, the correlation between the basal lattice parameter and the Fe site occupancy exhibited by the as-synthesised samples is maintained (Figure 11(a)). For the daughter of the non-superconducting parent sample (daughter B in Table 1), the combination of increasing the Fe1 site occupancy and the Li occupancy of the Li/Fe site by 13% results in the shortest a parameter of all the compounds investigated and a correspondingly large increase in c. These experiments also show that filling the iron site in the selenide layer, with its inevitable reduction of the Fe oxidation state, produces maximum values of Tc.
hydroxide layer moving to the site in the selenide layer 4.7 Å distant, presumably via the face of the Se4 tetrahedron forming the target site. This migration may be enabled by the fact that the metal hydroxide layer is relatively flat (Figure 10(b)) and the SCXRD measurements show that the Li/Fe ellipsoid is elongated along c.
Figure 10(a). Rietveld refinement against NPD data of the lithiated product depicted in Figure 9(a) (daughter A) with a refined composition Li0.84Fe0.16(OH)FeSe. The data from the 168° bank are displaced 4 units along the vertical axis. (b) Changes in refined positional parameters on lithiation of a superconducting parent material (parent A) to obtain daughter A. 50% displacement ellipsoids are shown for the lithiated daughter product; refined parameters for parent A are in parentheses. Further changes in refined parameters on lithiation are given in Table 1. The arrow shows a possible pathway for migration of iron. Table 1. Changes in lattice parameters and refined site occupancies on lithiation from NPD data. a (Å) Parent A 3.7893(2) Daughter A 3.7760(1) Parent B 3.8142(3) Daughter B 3.7542(1)
c (Å) 9.2617(6) 9.3512(2) 9.1882(7) 9.5859(3)
occ. Fe1 0.961(4) 1.004(5) 0.870(5) 1.000(8)
occ. Li 0.812(2) 0.837(2) 0.808(2) 0.934(8)
Total Fe 1.15(1) 1.665(5) 1.064(5) 1.07(1)
The reaction with Li/NH3 solution evidently provides additional Li which occupies the hydroxide layer at the expense of the Fe originally on that site. The Fe on the Li/Fe site is a reservoir and we presume that some of the iron migrates directly to the selenide layer to completely fill the site vacancies on the Fe1 site. The shortest migration pathway involves the Fe ion from the
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Figure 11. (a) Extension of the correlation between Fe and Li site occupancies and the basal lattice parameter, a, to include the lithiated samples (all results from NPD data). Deviation from linearity occurs only when large numbers of Fe vacancies are filled in the lithiation reaction and the Li content increases by a correspondingly large amount. Parent and daughter samples are linked for clarity. (b) The correlation between superconducting Tc and the average number of valence electrons per Li1−xFex(OH)Fe1−ySe formula unit assigned to the iron atoms in the selenide layers (a parameter that takes into account the iron deficiency in the selenide layer and the iron oxidation state).
Figure 11(b) shows a plot which includes all the samples made by the hydrothermal treatment and the post-synthetic lithiation. The
quantity plotted on the horizontal axis is the number of electrons associated with the iron atoms in the selenide layer per Li1– xFex(OH)Fe1–ySe formula unit (i.e. per selenide ion), making the assumption that the Fe ions present in the hydroxide layers are in the +2 oxidation state. This quantity takes into account both the site occupancy on the Fe1 site and the iron oxidation state. The plot of this quantity against Tc (defined to be 0 K if the compound showed no superconductivity) shows a smooth correlation of increasing Tc with increasing Fe electron count per formula unit once a threshold value is reached. These results are consistent with the phase diagrams found for related iron arsenide and selenide systems. For example: the electron count is a key controller in the iron arsenides and in the alkali metal iron selenide systems such as K0.8Fe1.6Se2 high site occupancies and iron oxidation states slightly below +2 found in portions of the samples are required for superconductivity. A further matter which requires further investigation is whether these hydroxide selenide samples show an inhomogeneous distribution of Fe1 site vacancies: the sample “Parent A” in Figure 9(a) shows evidence for a stepped magnetic susceptibility which could signify regions of the sample with different Tcs.
Measurement of superconducting properties Muon-spin rotation spectroscopy The lithiated sample depicted in Figure 10 (daughter A: refined composition Li0.84Fe0.16(OH)FeSe) was used for preliminary muon-spin rotation (μSR) spectroscopy measurements to probe the superconducting state in more detail. 300 mg of sample was contained in a silver foil packet and the sample was sequestered from air prior to loading into the helium atmosphere of the cryostat. Measurements were carried out on the MuSR beamline at the ISIS facility to measure the increase in Brms, the root-meansquare width of the magnetic field distribution, due to the development of the superconducting vortex lattice below Tc. The average field is also followed as a function of temperature and shows a diamagnetic response below Tc. Our results suggest a superconducting volume fraction above 50 %. Figure 12 shows Brms and –B0 as a function of temperature measured using a transverse field B0 of 10 mT and 30 mT. It is noticeable that Brms increases with B0 even in the normal state, which shows that there is a field-dependent contribution to the magnetic field distribution experienced by the muon that is most likely due to the presence of magnetism, and may be a consequence of the significant paramagnetic background originating from the Fe2+ ions in the hydroxide layer. In contrast, the diamagnetic response, measured by –B0 and which reflects the superconducting state only, is essentially unchanged between 10 and 30 mT. The different contributions to Brms act in quadrature (Brms2=Σ brms2), so using the proportionality between the superconducting contribution to Brms and 1/λab2, where λab=31/4λ is the in-plane penetration depth, we can extract an estimate of λab of 0.32(3) m from the 10 mT data, where the relatively large error takes account of the uncertainty due to field-induced effects. Our value of λab places this sample Li0.84Fe0.16(OH)FeSe close to the main scaling line in a Uemura plot of Tc against superfluid stiffness ρs=c2/λab2 (see inset).
9
Figure 12. The results of transverse-field muon-spin rotation spectroscopy on the lithiated sample “Daughter A” (see Figure 10; Table 1).
High field magnetometry Figure 13 shows the magnetisation as a function of magnetic field in a lithiated sample with Tc = 40 K, as determined by SQUID magnetometry. This shows characteristics of a type-II superconductor. Because the lower critical field Hc1 is very small, the Meissner effect is only apparent at lower temperatures, and there is a significant underlying paramagnetic behaviour, presumably arising from the Fe2+ moments in the hydroxide layer. The inset to Figure 13 shows the approximate evolution of the lower critical field with temperature, as deduced from the susceptibility, calculated from the magnetisation data in the main figure. A correction for the effect of the paramagnetic Fe2+centres in the hydroxide layer (Figure S9) yields no evidence for the upper critical field Hc2, so we may deduce that 0Hc2 > 7 T, which is in line with the behaviour of other iron-based superconductors.
Figure 13. Magnetisation as a function of magnetic field for a lithiated sample with Tc = 40 K. In the inset, open circles illustrate the field at which the calculated susceptibility is equal to zero, bars illustrate the approximate width in H of the transition from the Meissner state to the vortex lattice state, and the dashed line is a guide to the eye.
Conclusions Hydrothermal synthesis is a moderate temperature route to new iron-based superconductors in which iron selenide layers are separated by hydroxide layers. We describe a wide range of lithium iron hydroxide selenide compositions synthesised using a method similar to that described in refs 27 and 28, but using different source materials and no high pressure annealing treatment. The hydrothermal reactions produce compounds Li1−xFex(OH)Fe1−ySe where x ~ 0.2 and y ranges from 0.02 for reactions with larger Fe:Se ratios and reaches values as high as 0.15 for less iron-rich syntheses. The results of X-ray and neutron diffraction investigations coupled with magnetometry and X-ray absorption spectroscopy show that superconductivity in these compounds occurs when the iron deficiency in the selenide layers (i.e. y) is small which also implies low mean oxidation states for iron, and results in short Fe−Fe distances within the layers. The hydroxide layers separating the iron selenide layers fortuitously contain a source of further iron atoms, which, in a post-synthetic reductive lithiation treatment can be displaced by lithium ions to completely fill the vacant sites in the iron selenide layers (y becomes 0 within the experimental uncertainty). This enables the full range of hydrothermally as-synthesised samples to be transformed into superconductors with Tc exceeding 40 K. These results quantify the correlation between the occurrence of superconductivity and the size of Tc with the composition and structure of the iron selenide layer and electron count in a recently discovered series of iron-based superconductors.
Acknowledgements This work was funded by the Engineering and Physical Sciences Research Council (grant EP/I017844) and the Leverhulme Trust (grant RPG-2014-221). We also thank the Diamond Light Source Ltd. for the award of beamtime (MT9981 on I19; EE9697 on I11; SP11061 on B18) and studentship support for SJC (Cassidy), and the ISIS facility for the award of neutron beamtime including awards under the GEM Xpress programme. SJS acknowledges the support of a DFG Fellowship (SE2324/1-1). We are grateful for the assistance of beamline scientists at Diamond (G. Cibin, C. C. Tang) and ISIS (R. I. Smith, P. J. Baker), Dr N. Charnley, Oxford Earth Sciences, for assistance with SEM measurements and Dr R. Jacobs for assistance with TGA measurements.
Notes and references a
Department of Chemistry, University of Oxford, Inorganic Chemistry Laboratory, South Parks Road, Oxford, OX1 3QR, UK. Fax: 44 1865 272690; Tel: 44 1865 272695; E-mail:
[email protected] b Beijing Synchrotron Radiation Facility, Institute of High Energy Physics, Chinese Academy of Science, Beijing 100049, China. c Diamond Light Source Ltd., Harwell Science and Innovation Campus, Didcot, OX11 0DE, UK. d Materials and Engineering Research Institute, Faculty of Arts, Computing, Engineering and Sciences, Sheffield Hallam University, City Campus, Howard Street, Sheffield, S1 1WB,UK e School of Physical Sciences, Ingram Building, University of Kent, Canterbury, Kent, CT2 7NH, UK. f Department of Physics, University of Oxford, Clarendon Laboratory, Parks Road, Oxford, OX1 3PU, UK.
10
† Electronic Supplementary Information (ESI) appended: Further Rietveld refinements, tables of refined parameters, Crystallographic Information File for single crystal refinements, further X-ray absorption data and fits, further magnetometry analysis. ‡ Li1–xFex(OH)Fe1–ySe: Space group P4/nmm (No. 129) a ~ 3.8 A, c ~ 9.2 Å. Atomic positions (origin choice 2: inversion centre at origin): Fe: site 2a (¼, ¾, 0); Se site 2c (¾, ¾, z~0.16); O site 2c (¼, ¼, z~0.43); Li/Fe site 2b (¾, ¼, ½); H site 2c (¼, ¼, z~0.33). Full details in tables S1 and S2 and in the crystallographic information file available on request.
1
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24
M. Burrard-Lucas, D. G. Free, S. J. Sedlmaier, J. D. Wright, S. J. Cassidy, Y. Hara, A. J. Corkett, T. Lancaster, P. J. Baker, S. J. Blundell and S. J. Clarke, Nat. Mater., 2013, 12, 15. 25 E.-W. Scheidt, V. R. Hathawar, D. Schmitz, A. Dunbar, W. Scherer, F. Mayr, V. Tsurkan, J. Deisenhofer, and A. Loidl, Eur. Phys. J. B., 2012, 85, 279. 26 S. J. Sedlmaier, S. J. Cassidy, R. Morris, M. Drakopoulos, C. Reinhard, S. J. Moorhouse, D. O'Hare, P. Manuel, D. Khalyavin, and S. J. Clarke, J. Am. Chem. Soc., 2014, 136, 630. 27 X. F. Lu, N. Z. Wang, G. H. Zhang, X. G. Luo, Z. M. Ma, B. Lei, F. Q. Huang, and X. H. Chen, Phys. Rev. B, 2013, 89, 020507(R). 28 X. F. Lu, N. Z. Wang, H. Wu, Y. P. Wu, D. Zhao, X. Z. Zeng, X. G. Luo, T. Wu, W. Bao, G. H. Zhang, F. Q. Huang, Q. Z. Huang, X. H. Chen, arXiv:1408.2006. 29 P. F. Ziemkiewicz, M. O’Neal and R. J. Lovett, Mine Water Environ., 2011, 30, 141. 30 G. A. Kakos, T. W. Turney and T. B. Williams, J. Solid State Chem. 1994, 108, 102. 31 S. J. Blundell, Magnetism in Condensed Matter, Oxford University Press, 2001. 32 S. P. Thompson, J. E. Parker, J. Potter, T. P. Hill, A. Birt, T. M. Cobb, F. Yuan and C. C. Tang, Rev. Sci. Instrum., 2009, 80, 075107. 33 H. Nowell, S. A. Barnett, K. E. Christensen, S. J. Teat, and D. R. Allan, J. Synchrotron Radiation, 2012, 19, 435. 34 P. W. Betteridge, J. R. Carruthers, R. I. Cooper, K. Prout and D. J. Watkin, J. Appl. Cryst., 2003, 36, 1487. 35 A. A. Coelho, TOPAS Academic, Version 5; Coelho Software, Brisbane, Australia, 2012. 36 L. Palatinus and G. Chapuis, J. Appl. Crystallogr., 2007, 40, 786. 37 T. Minagawa, Memoirs of Osaka Kyoiku University, Series III, 1994, 43, 39. 38 V. F. Sears, Neutron News, 1992, 3, 29. 39 V. G. Young, jr., M. J. McKelvy, W. S. Glaunsinger and R. B. von Dreele, Solid State Ionics, 1988, 26, 47. 40 M. J. Pitcher, D.Phil. Thesis, University of Oxford. 2011. 41 P. Adamson, J. Hadermann, C. F. Smura, O. J. Rutt, G. Hyett, D. G. Free and S. J. Clarke, Chem. Mater., 2012, 24, 2802. 42 J. K. M Lindén, E.-L. Rautama, M. Karppinen and H. Yamauchi, Hyperfine Interactions 2012, 208, 133. 43 A. G. Maddock. Mössbauer Spectroscopy. Horwood Publishing, 1997, 108. 44 A. J. Dent, G. Cibin, S. Ramos, A. D. Smith, S. M. Scott, L. Varandas, M. R. Pearson, N. A. Krumpa, C. P. Jones, P. E. Robbins, J. Phys. Conf. Ser. 2009, 190, 012039. 45 B. Ravel and M. Newville, J. Synchrotron Radiat., 2005, 12, 537. 46 S. J. Cassidy, S. Ramos, and S. J. Clarke, Z. Anorg. Allg. Chem. 2014, (Submitted). 47 A. Iadecola, B. Joseph, L. Simonelli, A. Puri, Y. Mizuguchi, H. Takeya, Y. Takano and N. L. Saini, J. Phys.: Condens. Matter 2012, 24, 115701. 48 T. A. Tyson, T. Yu, S. J. Han, M. Croft, G. D. Gu, I. K. Dimitrov and Q. Li, Phys. Rev. B, 2012, 85, 024504.
11
Electronic Supplementary Information Summary of refined structural parameters Table S1. Refined structural parameters from Neutron Powder Diffraction measurements. Sample
shl4p1
Comment
shl18p3
shl18p3
SHL13p4
shl27p2
shl27p4
shl26p2
Parent B
Tc (K)
N/A
N/A
Instrument
GEM
GEM
T (K)
295
295
a (Å)
3.82101
c (Å)
shl17p3
shl17p3
shl28p4
GMA128
GMA127
Daughter A
Daughter B
39
42
41
Parent A N/A
N/A
17
23
26
GEM
HRPD
HRPD
HRPD
HRPD
GEM
GEM
GEM
HRPD
HRPD
50
295
295
295
295
295
50
295
295
295
3.8142(3)
3.8021(1)
3.81658
3.80395(3)
3.79330(3)
3.78943(4)
3.7893(2)
3.7757(1)
3.790(2)
3.77599(7)
3.7543(1)
9.16377
9.1882(7)
9.0885(3)
9.17525
9.20892(9)
9.24734(9)
9.2583(1)
9.2616(6)
9.1248(3)
9.2486(6)
9.3512(2)
9.5859(3)
3
V (Å )
133.7921
133.67(2)
131.382(9)
133.6493
133.254(2)
133.061(3)
132.948(3)
132.99(2)
130.080(8)
132.85(2)
133.331(6)
135.108(9)
c/a
2.398258
2.4089(3)
2.3904(1)
2.40405
2.42089(3)
2.43781(3)
2.44320(4)
2.4442(2)
2.4167(1)
2.4403(2)
2.47648(7)
2.5534(1)
occ(Fe1)
0.846(5)
0.870(5)
0.876(4)
0.880(4)
0.921(4)
0.950(4)
0.956(4)
0.961(4)
0.955(4)
0.972(4)
1.004(5)
1.000(8)
occ(Li)
0.800(2)
0.805(2)
0.802(2)
0.802(1)
0.811(1)
0.816(1)
0.816(1)
0.812(2)
0.808(1)
0.807(2)
0.837(2)
0.930(7)
occ(H)
1.01(1)
1*
1.000(8)
0.964(4)
0.970(8)
0.950(8)
0.903(9)
1.02(1)
0.963(8)
0.990(9)
0.95(1)
0.98(2)
Total Fe
1.046(7)
1.065(7)
1.074(6)
1.078(5)
1.110(5)
1.134(5)
1.140(5)
1.149(6)
1.147(5)
1.165(6)
1.163(7)
1.07(1)
zSe
0.1613(2)
0.1626(2)
0.1623(2)
0.1623(1)
0.1621(1)
0.1619(1)
0.1622(1)
0.1625(2)
0.1624(1)
0.1632(2)
0.1604(2)
0.1573(2)
zO
0.4280(3)
0.4270(4)
0.4265(3)
0.4273(2)
0.4263(2)
0.4263(2)
0.4258(2)
0.4253(3)
0.4256(2)
0.4257(3)
0.4246(3)
0.4229(4)
zH
0.3301(7)
0.3320(7)
0.3253(5)
0.3283(5)
0.3261(4)
0.3273(4)
0.3266(4)
0.3267(5)
0.3236(4)
0.3257(5)
0.3292(6)
0.3285(9)
Fe–Se (Å)
2.416(1)
2.419(1)
2.4064(9)
2.4203(6)
2.4179(5)
2.4165(6)
2.4176(7)
2.420(1)
2.3998(8)
2.423(1)
2.411(1)
2.408(1)
Se–Fe–Se ×2 () (°)
104.54(7)
103.84(7)
104.37(5)
104.08(4)
103.74(3)
103.42(3)
103.21(4)
103.07(7)
103.75(5)
102.92(6)
103.08(5)
102.46(7)
Se–Fe–Se ×4 () (°)
111.99(4)
112.36(4)
112.08(3)
112.23(2)
112.41(2)
112.58(2)
112.69(2)
112.76(4)
112.40(3)
112.84(3)
112.76(3)
113.09(4)
2
0.0157(7)
0.0188(7)
0.0066(4)
0.0191(4)
0.0182(3)
0.0161(3)
0.0161(3)
0.0115(4)
0.0043(2)
0.0114(4)
0.0145(5)
0.0157(7)
2
0.0197(7)
0.0261(8)
0.0079(4)
0.0192(4)
0.0193(4)
0.0181(4)
0.0194(4)
0.0148(5)
0.0071(3)
0.0157(5)
0.0178(5)
0.0157(8)
2
0.018(1)
0.018(1)
0.0098(6)
0.0201(6)
0.0214(5)
0.0226(5)
0.0232(6)
0.0186(8)
0.0091(5)
0.0194(8)
0.0206(8)
0.019(1)
Ueq (H) (Å )
2
0.053(2)
0.062(3)
0.035(2)
0.053(1)
0.054(1)
0.053(1)
0.050(1)
0.055(2)
0.035(1)
0.058(2)
0.049(2)
0.054(3)
wRp
0.0301
0.0442
0.0336
0.0285
0.0282
0.0305
0.0301
0.0299
0.0261
0.0335
0.0245
0.0292
Ueq (Fe) (Å ) Ueq (Se) (Å ) Ueq (O) (Å )
P4/nmm (No. 129) atomic positions (origin choice 2: inversion centre at origin): Fe 2a (¼, ¾, 0); Se 2c (¾, ¾, zSe); O 2c (¼, ¼, zO); Li/Fe 2b (¾, ¼, ½); H 2c (¼, ¼, zH) * not refined
S1
Table S2. Summary of refinement results at 100 K against I19 single crystal X-ray diffraction data. The higher precision, and more reproducible lattice parameters obtained at room temperature using model-independent Pawley fits to powder diffraction data obtained on I11 and/or Rietveld refinements against room temperature NPD data are also included and are the values used for the plots in the main text. Sample
shl4p1
shl18p3
shl13p2
shl12p1
shl11p2
Tc (K)
0
0
0
0
16
shl11p4† 20
shl12p3
shl17p3
shl9p3
shl5p2
shl28p4
23
26
30*
37
39
Powder diffraction Instrument
I11/GEM
GEM
I11
I11
I11
I11
I11
GEM
I11
I11
GEM
T (K)
295
295
295
295
295
295
295
295
295
295
295
a (Å)
3.82101(1)
3.8142(3)
3.81736(1)
3.79913(1)
3.78998(1)
3.78882(1)
3.78759(1)
3.7893(2)
3.78719(1)
3.78500(1)
3.7900(2)
c (Å)
9.16377(4)
9.1882(7)
9.17522(1)
9.22470(1)
9.26172(2)
9.26564(5)
9.26126(1)
9.2616(6)
9.26379(6)
9.26482(5)
9.2486(6)
c/a
2.39826(1)
2.4089(3)
2.40355(1)
2.42811(1)
2.44374(1)
2.44552(2)
2.44516(1)
2.4442(2)
2.44609(2)
2.44777(2)
2.4402(2)
3
133.792(1)
133.67(2)
133.704(1)
133.144(1)
133.035(1)
133.010(1)
132.861(1)
132.99(2)
132.869(2)
132.730(1)
132.85(2)
V (Å )
Single crystal X-ray diffraction Instrument
I19
I19
I19
I19
I19
I19
I19
I19
I19
I19
I19
I19
Wavelength (Å)
0.6889
0.6889
0.6889
0.6889
0.6889
0.6889
0.6889
0.6889
0.6889
0.6889
0.6889
0.6889
T (K)
100
100
100
100
100
100
100
100
100
100
100
100
a (Å)
3.8219(5)
3.8020(1)
3.8038(1)
3.8249(1)
3.8025(1)
3.79441(8)
3.7861(2)
3.7958(1)
3.7782(1)
3.7709(2)
3.7793(4)
3.7677(2)
c (Å)
9.175(2)
9.1098(2)
9.1052(5)
9.2378(7)
9.2345(5)
9.2171(5)
9.2228(7)
9.2207(7)
9.1728(4)
9.1562(6)
9.192(2)
9.155(1)
3
V (Å )
134.01(4)
131.684(4)
131.740(9)
135.15(1)
133.523(8)
132.703(8)
132.21(1)
132.86(1)
130.941(8)
130.20(1)
131.29(4)
129.96(2)
occ. Fe1
0.87(3)
0.875(7)
0.860(5)
0.920(8)
0.960(16)
0.97(1)
0.974(5)
0.973(6)
0.972(6)
0.982(11)
0.958(7)
0.982(7)
occ. Li on Li/Fe site occ. Fe on Li/Fe site Fe:Se ratio
0.81(4)
0.811(8)
0.800(6)
0.824(10)
0.81(2)
0.811(13)
0.829(6)
0.823(7)
0.801(7)
0.789(15)
0.810(8)
0.794(10)
0.19(4)
0.189(8)
0.200(6)
0.176(10)
0.19(2)
0.189(13)
0.171(6)
0.177(7)
0.199(7)
0.211(15)
0.190(8)
0.206(10)
1.06(7)
1.064(15)
1.060(11)
1.096(18)
1.15(4)
1.16(2)
1.15(1)
1.150(12)
1.171(13)
1.19(3))
1.148(15)
1.188(18)
zSe
0.1619(3)
0.16175(8)
0.16208(8)
0.16135(12)
0.16075(16)
0.16097(12)
0.16106(5)
0.16116(7)
0.16116(6)
0.16097(11)
0.16088(9)
0.16139(9)
zO
0.427(3)
0.4285(6)
0.4285(6)
0.4273(10)
0.4288(17)
0.4281(12)
0.4275(5)
0.4265(5)
0.4273(5)
0.4275(10)
0.4280(7)
0.4269(8)
Fe-Se (Å) × 4
2.4206(18)
2.4052(4)
2.4073(4)
2.4247(7)
2.4121(9)
2.4085(7)
2.4060(4)
2.4105(4)
2.3988(3)
2.3932(6)
2.3995(6)
2.3942(5)
Se–Fe–Se ×2 () (°) 2 Ueq Fe1 (Å )
104.27(11)
104.44(3)
104.38(3)
104.14(4)
104.04(6)
103.95(4)
103.76(2)
103.88(2)
103.91(2)
103.97(4)
103.91(3)
103.79(3)
0.031(2)
0.0095(5)
0.0149(5)
0.0153(9)
0.0127(11)
0.0108(8)
0.0109(5)
0.0087(5)
0.0083(4)
0.0092(7)
0.0121(7)
0.0107(6)
2
0.029(3)
0.0094(4)
0.0146(3)
0.0156(5)
0.0124(9)
0.0106(6)
0.0108(3)
0.0094(3)
0.0080(3)
0.0091(7)
0.0127(4)
0.0109(4)
2
0.035(5)
0.012(3)
0.018(2)
0.022(4)
0.019(5)
0.018(4)
0.019(2)
0.015(2)
0.014(2)
0.017(4)
0.019(3)
0.019(3)
2
0.0554
0.026(3)
0.036(3)
0.035(5)
0.034(6)
0.032(5)
0.028(2)
0.027(2)
0.030(2)
0.033(6)
0.0336
0.036(3)
Ueq Se(Å ) Ueq O (Å ) Ueq Li/Fe (Å )
S2
2
0.0218(19)
0.0070(4)
0.0108(4)
0.0111(6)
0.0092(10)
0.0063(7)
0.0071(3)
0.0042(3)
0.0061(3)
0.0072(6)
0.0067(4)
0.0066(5)
2
0.050(3)
0.0143(6)
0.0230(5)
0.0236(9)
0.0196(12)
0.0199(9)
0.0183(4)
0.0177(5)
0.0126(4)
0.0133(8)
0.0229(7)
0.0190(6)
2
0.0211(11)
0.0086(3)
0.0123(2)
0.0127(4)
0.0103(8)
0.0072(5)
0.0082(2)
0.0059(2)
0.0064(2)
0.0074(3)
0.0084(2)
0.0076(3)
2
0.0452(17)
0.0111(4)
0.0194(3)
0.0215(6)
0.0166(9)
0.0175(6)
0.0160(3)
0.0165(3)
0.0114(3)
0.0125(5)
0.0214(4)
0.0175(4)
2
0.027(6)
0.0096(16)
0.0130(15)
0.018(3)
0.010(4)
0.012(3)
0.0139(11)
0.0094(14)
0.0110(11)
0.014(2)
0.015(2)
0.014(2)
2
0.052(11)
0.018(3)
0.026(3)
0.029(4)
0.035(7)
0.031(5)
0.029(2)
0.026(3)
0.020(2)
0.024(5)
0.027(3)
0.028(3)
2
0.020(8)
0.0088(18)
0.0104(16)
0.013(3)
0.012(4)
0.013(3)
0.0104(14)
0.0091(16)
0.0119(14)
0.018(3)
0.014(2)
0.016(2)
U33 Li/Fe (Å )
2
0.13(2)
0.060(4)
0.090(5)
0.078(7)
0.078(12)
0.069(7)
0.062(3)
0.064(4)
0.065(4)
0.063(7)
0.073(5)
0.076(6)
R1 (all)
0.1119
0.0204
0.025
0.0403
0.0492
0.0364
0.0267
0.0362
0.0332
0.0546
0.055
0.0419
wR (all)
0.216
0.0668
0.0467
0.0749
0.2165
0.1286
0.0557
0.0741
0.0677
0.1358
0.081
0.0804
R1 (I >2(I))
0.0889
0.0198
0.0197
0.0302
0.0457
0.0342
0.0202
0.0317
0.0229
0.0393
0.0344
0.0336
wR (I >2(I))
0.2033
0.0664
0.0458
0.0741
0.1941
0.1273
0.0456
0.0709
0.0648
0.129
0.0733
0.0774
Sample
shl4p1
shl18p3
shl13p2
shl12p1
shl11p2
shl11p4
shl11p4
shl12p3
shl17p3
shl9p3
shl5p2
shl28p4
I19 Collection serial number
014DNW14
010DNW14
020DNW14
021DNW14
017DNW14
018DNW14
003DNW14
001SJS14
013DNW14-
016DNW14
012DNW14
025DNW14
U11 Fe1 (Å ) U33 Fe1 (Å ) U11 Se (Å ) U33 Se (Å ) U11 O (Å ) U33 O (Å ) U11 Li/Fe (Å )
P4/nmm (No. 129) atomic positions (origin choice 2: inversion centre at origin): Fe 2a (¼, ¾, 0); Se 2c (¾, ¾, zSe); O 2c (¼, ¼, zO); Li/Fe 2b (¾, ¼, ½); H 2c (¼, ¼, zH). For the SCXRD data collections the H atom was included with the displacement parameter fixed to be 2.5 times that of O and the O–H distance restrained to be the 0.91 Å obtained from the NPD refinements. * evidence for superconducting transitions at 30 K and 22 K. † Two different crystals were measured from one batch and gave consistent results.
S3
Rietveld Refinements against neutron powder diffraction data Figure S1. Refinements against Neutron Powder Diffraction Data (HRPD and GEM instruments at ISIS) measured at room temperature for hydrothermally synthesised samples. Data (black line), fit (red line) and difference (grey line) are shown together with reflection markers for the phases included in the refinement. For small samples the vanadium container was included in the fit as an additional phase. For multiphase samples compositions are expressed as weight %.
S4
Figure S2. (a)–(d) Refinements against Neutron Powder Diffraction Data (GEM instruments at ISIS; BANK 2 at 2 = 18° shown) measured at room temperature and 50 K for two hydrothermally synthesised samples (SHL18p3 (a) and (b) (Parent B in the main text) and SHL17p3 (c) and (d) (Parent A) subsequently used for lithiation). Data (black line), fit (red line) a nd difference (grey line) are shown together with reflection markers for the phases included in the refinement. For small samples the vanadium container was included in the fit as an additional phase. The presence of temperature-invariant broad features in the diffraction pattern (*) for the superconductor SHL17p3 are also observed in SHL28p4 (e), and these may be intrinsic to these iron -rich phases.
S5
Figure S3. Refinements against Neutron Powder Diffraction Data (HRPD instruments at ISIS) measured at room temperature for samples lithiated after hydrothermal synthesis. Data (black line), fit (red line) and difference (grey line) are shown together with reflection markers for the phases included in the refinement. The vanadium container was included in the fit as an additional phase. Compositions are expressed as weight %. Sample GMA127 (Daughter B) was derived from sample SHL18p3 (Parent B) (Figure S1(f)) and a small residue of this material remained after lithiation. Sample GMA128 (Daughter A) was derived from SHL17p3 (Parent A) (Figure S1(g)).
S6
Effects of air exposure on superconducting and magnetic properties. Figure S4. (a) Evolution of the magnetic susceptibility with exposure to ambient air of a superconducting sample obtained directly from hydrothermal synthesis (SHL28p4). (b) Fits to the molar magnetic susceptibility obtained in measurements conducted in the paramagnetic region for the same fresh and 1-week-air-exposed sample as in (a). These measurements were made by measuring the magnetic moment of the sample as a function of temperature at applied fields of 4T and 3T as the magnetisation was found to vary linearly with field at these fields, but showed some deviation at low fields due to the presence of minuscule amounts of ferromagnetic impurity below the detection limits of our diffraction experiments . The table gives a 2+ value for the effective moment (eff) calculated per Fe ion in the hydroxide layer which is consistent with Fe in the as made 3+
material being oxidised to Fe in the air-exposed product.
S7
Figure S4(c). The effect of air exposure and annealing under vacuum at the synthesis temperature on a hydrothermally synthesised sample (sample SHL12p3 with composition Li 0.82 Fe0.18(OH)Fe0.97 Se from SCXRD). Note that only a few points were measured after field-cooling of the annealed sample, sufficient to demonstrate the Meissner effect (red open circles).
Figure S4(d). The effect of air exposure (10 days) on the lithiated sample (Daughter A) shown in Figure 9(a)
S8
SEM analysis. Figure S5. SEM micrographs of typical hydrothermally-synthesised samples (Parent A – left; Parent B – right). The crystals have a thickness of about 1 m.
S9
Thermogravimetric analysis. Figure S6. Data collected under flowing N2 for the non-superconducting sample SHL18p3 (parent B) (black) show a mass loss of 5.4% consistent with non-redox dehydration to maintain a mean Fe oxidation state of +2.08: Li 0.8 Fe1.07 Se(OH)
→
Li 0.8 Fe1.07 SeO 0.5 + ½ H2 O
Data for the more iron-rich superconductor SHL17p3 (Parent A) show a mass loss of 4.6 % which could be accounted for by dehydration coupled with mild oxidation from the environment to produce a product with a similar mean Fe oxidation state as in the decomposition product of SHL18p3. i.e. the mass loss of 4.6% for Li 0.8 Fe1.15 Se(OH) would be accounted for by a product Li 0.8 Fe1.15 SeO 0.59 It was not possible to reliably determine the identity of the products formed on oxidation from powder diffraction measurements.
S10
Analysis of EXAFS data It is observed in the EXAFS that vacancies in the [Fe1-ySe] layer lead to a loss of local ordering outside of the first coordination shell. For samples with high vacancy concentrations there is no appreciable EXAFS signal beyond the first coordination shell (R>3). For samples with l ow vacancy concentrations a good fit to the full EXAFS can be achieved out to a radial distance of 5.5 Å from the absorber by modelling single and multiple scattering paths originating entirely from the FeSe layer, with the addition of the Fe–O path from the Li/Fe(OH) layer which contributes towards the first coordination shell. These two cases are demonstrated in Figure S7(a) and (b) for a high-vacancy-containing, hydrothermally synthesised material Li 0.80 Fe0.20(OH)Fe0.88Se (similar by diffraction investigations to the sample designated Parent B in the main article) and the vacancy free product obtained by lithiation Li 0.93Fe0.07 (OH)Fe1.00 Se (Daughter B). An equivalent fit to hydrothermally synthesised materials with low vacancy concentrations can be achieved using the same model as for the lithiated materials, shown for Li 0.81 Fe0.19 (OH)Fe0.96Se (Parent A) in Figure S7(c).
(a)
(b)
(c)
Fi gure S7 (a ) fi ts to the Fe K-edge EXAFS for an iron-poor, non-superconducting sample Li 0.80Fe0.20(OH)Fe0.88Se and (b) the lithiated s uperconducting product Li 0.93Fe0.07(OH)Fe1.00Se. (c) The fit to the Fe K-edge EXAFS for an iron-rich superconducting hydrothermallys ynthesised compound, Li 0.81Fe0.19(OH)Fe0.96Se
S11
Table S3. Refinements show strong agreement between the average structure as probed by diffraction techniques and the local structure of Li 1-xFex(OH)Fe1-ySe with low values of y. Path
st
Fe1-Se 1 neighbour, single scattering st Fe-Fe 1 neighbour, single scattering st Fe2-O 1 neighbour, single scattering st Fe1-Se-Fe 1 neighbours, acute triangle nd Fe1-Fe 2 neighbour, single scattering st nd Fe1-Se-Fe 1 and 2 neighbours, double scattering nd Fe1-Se 2 neighbour, single scattering rd Fe-Fe 3 neighbour, single scattering and forward st scattering paths through 1 neighbouring Fe
Li 0.80 Fe0.20 (OH)Fe0.88 Se (see Figure S7(a)) difference between EXAFS and diffraction calculated path length (Å) 0.003(4)
Li 0.81 Fe0.19 (OH)Fe0.96 Se (see Figure S7(c)) difference between EXAFS and diffraction calculated path length (Å)
Li 0.93 Fe0.07 (OH)Fe1.00 Se (see Figure S7(b) difference between EXAFS and diffraction calculated path length (Å)
0.003(3)
0.007(4)
0.027(9)
-0.003(4)
-0.017(6)
0.00(3)
-0.02(2)
0.05(6)
Not modelled
0.04(6)
-0.03(4)
Not modelled
0.01(6)
-0.01(3)
Not modelled
0.0(1)
0.02(7)
Not modelled
0.03(6)
-0.02(3)
Not modelled
0.03(2)
-0.09(3)
The ratio of iron in the Fe1-ySe layer and the Li 1-xFex(OH) layer was refined from the EXAFS data by comparing relative contributions to the EXAFS from the Fe1 site and Fe in the Li/Fe layer. Since the only significant contribution to the Fe EXAFS originating from the Li/Fe(OH) layer is the Fe-O path in the first coordination shell, it is was appropriate to extract this ratio from this region alone so as not to introduce unnecessary correlations. The same model was applied for all materials consisting of an Fe1 -Se path from the Fe1-ySe layer, an Fe-O path from the Li 1xFex(OH) layer, and an Fe-Fe path which is present in both layers the distance of which is 1/√2 times the a unit cell parameter. The Fe-O path was fixed at its crystallographic value as determined by diffraction due to its relatively low contribution to the EXAFS the X-ray data gives a more reliable distance. All other path lengths and Debye-Waller terms were allowed to refine 2 freely. The ratio of iron between the two layers was refined by multiplying the amplitude reduction factor (S 0 ) of the Fe-O path by a weighting factor w. Fits for 6 materials with varying iron vacancy concentrations are shown in Figure S8, the relative contributions of the Fe-O and Fe1-Se paths are found to be consistent with the relative concentrations of iron in each layer refined by diffraction methods (see Table S4).
S12
Li 0.80 Fe 0.20 (OH)Fe0.88Se
Li 0.81 Fe 0.19 (OH)Fe0.92 Se
Li 0.81 Fe 0.19 (OH)Fe0.96Se
Li 0.82 Fe 0.18 (OH)Fe0.97 Se
Li 0.84 Fe 0.16 (OH)Fe1.00Se
Li 0.93 Fe 0.07 (OH)Fe1.00 Se
Fi gure S8. Fi ts to the first shell of the Fe K-edge EXAFS for Li 1-xFe x(OH)Fe1-y Se with varyi ng x a nd y. relative contributions from the three paths modelled are plotted below the fit. The relative contribution of the Fe -O path from Fe i n the Li/Fe site decreases with decreasing x a nd y.
Table S4. Comparison of the ratios of the occupancies of the two iron sites from EXAFS and diffraction methods. 2
Sample
Refined value from diffraction methods Refined value from the EXAFS for S0 for the ratio of iron in Li/Fe layer: FeSe the ratio of iron in Li/Fe layer: layer FeSe layer (w) Li 0.80 Fe0.20 (OH)Fe0.88 Se 0.23 0.26(6) 0.64(5) Li 0.81 Fe0.19 (OH)Fe0.92 Se 0.21 0.16(5) 0.72(5) Li 0.81 Fe0.19 (OH)Fe0.96 Se 0.20 0.16(3) 0.77(4) Li 0.82 Fe0.18 (OH)Fe0.97 Se 0.19 0.18(4) 0.78(5) Li 0.84 Fe0.16 (OH)Fe1.00 Se 0.16 0.08(3) 0.81(5) Li 0.93 Fe0.07 (OH)Fe1.00 Se 0.07 0.04(4) 0.79(5) This model approximates the zero point shift in the edge position (ΔE 0 ) to the same value for both iron sites. There is no observable splitting in the edge and correlations between shift and magnitude parameters are low so this assumption will have 2 negligible impact on the result. A single value is used to describe the amplitude reduction factor (S 0 ) for both Fe sites, while w describes the additional reduction in amplitude for Fe in the Li/Fe site arising from there being fewer sites, the largest 2
correlation between S0 and w was 0.46.
S13
High field magnetometry.
Figure S9. The magnetisation of the lithiated sample shown in Figure 13 (Tc = 40 K) with a paramagnetic contribution from the 2+
Fe ions in the hydroxide layer subtracted. From this, we see no evidence of Hc2 below 7 T. The subtraction was carried out in the following way: from the magnetisation data shown in Figure 13, one can calculate the susceptibility of the sample as a function of temperature and magnetic field. If we assume that at the highest temperature a nd 2+ field measured (T = 41 K and 0 H = 7 T) the superconductivity is completely suppressed and Fe spin-only paramagnetism (S = 2) is the sole contribution to the susceptibility, we can calculate the hypothetical number density of paramagnetic ions. Since, at this temperature and field, gSB0 H/kT = 0.46 < 1, it is valid to use Curie's law as an approximation of the full Brillouin function to describe paramagnetism. Thus 2
2
= n 0 B g S(S+1)/3kT 2+
where n is the number density of paramagnetic Fe ions (g = 2; S =2) 2+
For the sample in Figure 13, this gives 0.17 paramagnetic Fe ions per formula unit Li 1–xFex(OH)FeSe. This sample was similar, but not identical, to the sample Daughter A used in the SR measurements. n thus equates closely to the value of x in the formula which is in line with the results reported in Figure S4(a). With this information, we subtracted the expected paramagnetic magnetisation from the measured magnetisation to obtain the `reduced' magnetisation plo tted in Figure S9. This does not appear to show any feature indicative of the upper critical field, Hc2 . From this, we deduce that Hc2 > 7 T, which is in line with reports on many iron-based superconductors.
S14
