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FEATURE ARTICLE Design strategies in mineralized biological materials Stephen Weiner and Lia Addadi Dept. of Structural Biology, Weizmann Institute of Science, 76100 Rehovot, Israel Organisms have been producing mineralized skeletons for the past 550 million years. They have evolved many dierent strategies for improving these materials at almost all hierarchical levels from A° ngstroms to millimetres.Key components of biological materials are the macromolecules, which are intimately involved in controlling nucleation, growth, shaping and adapting mechanical properties of the mineral phase to function. One interesting tendency that we have noted is that organisms have developed several strategies to produce materials that have more isotropic properties. Much can still be learned from studying the principles of structure–function relations of biological materials.Some of this information may also provide new ideas for improved design of synthetic materials. Many biological materials are known to have unusual mechan- there are infrequent occurrences of preserved non-mineralized ical properties, some of which are surprisingly advantageous, fossils in much older rocks.5 The Cambrian fossil record, especially when taking into account the fact that they are however, does provide us with good documentation of the formed at ambient temperatures and pressures.1–3 This obser- very beginnings of skeletal evolution, and the remaining 500 vation has inspired many studies of these materials over the million years can be viewed, from our perspective, as an last several decades aimed at discovering some of their struc- ‘extensive’ product testing period.The evolution of biomineraltural ‘secrets’. The mineralized biological materials represent ization processes has been reviewed by Lowenstam and an interesting subgroup within this vast world. Clearly the Weiner.6 presence of the mineral phase and the manner in which the Members of many dierent phyla began forming mineralized mineral and the organic material are organized are among exoskeletons within several millions of years (at the base of the key factors that contribute to their unique mechanical the Cambrian).6,7 This followed what was probably a very properties.major extinction,8 and in fact the phylogenetic groups that Understanding just how this occurs is, however, not a trivial subsequently mineralized were themselves, for the most part, task, as the scale of ordered structures can vary from A° ngstroms newly evolved.From this we can infer that there was probably to millimetres. Furthermore, one level of structural organiz- some external ecological pressure on these organisms to ation is often nested into another larger-scaled structural type, develop mineralized materials (protection against predators is to produce a very complicated overall structure.4 In order to a favourite explanation),9 and that the genetic ‘backgrounds’ reveal the design strategies of these natural materials we not of the mineralizers were still rather flexible and hence amenable only have to understand their structural features in great detail, to novel experiments.We also know that today all the more but we also need to identify the specific benefits that a sophisticated mineralizers use similar families of proteins for particular aspect of the structure contributes to the bulk controlling mineral nucleation and growth, implying that material. With all these diculties, it is therefore not surprising underlying mechanisms common to many taxonomic groups that we still really know very little about the strategies used do exist.10 We do not, however, know if this capability was by organisms to form their superior materials.divergently inherited from a common ancestor or convergently In this review we will address key issues relating to min- arrived at independently by each group.eralized biological material formation and function at dierent These first skeletal mineralizers did not all use the same dimensional scales. We will start at the A° ngstrom level and mineral. About a third opted for a calcium phosphate mineral work our way up to the millimetre level. Our strategy will be (carbonated apatite) and most for a calcium carbonate mineral.to illustrate the concepts presented by focusing mainly on one Of the latter, almost all chose calcite from among the family family of important biogenic minerals, the calcium carbonates, of CaCO3 polymorphs. One phylum deposited amorphous and through them identify more general principles of biological silica (also known as opal) and one group of bacteria, magnet- materials design.We also discuss non-carbonate containing ite.6,11 Presumably other minerals were also used, but these materials, such as bone, where appropriate. We will begin, have not yet been discovered or were not preserved. So however, with more general information about mineralized environmental conditions at the time, such as the sea water biological materials. chemistry, did not apparently aect their options.Nor is there any reason to believe that the capabilities of these early mineralizers were limited. They formed an incredible array of Materials with a history morphologically varied structures, with a variety of minerals and presumably macromolecules as well.12 Many endo- and exo-skeletons are composed of mineralized One enigmatic observation emanating from the Cambrian materials.They have one unusual advantage over other biologifossil record is that the microstructures of many of the min- cal and synthetic materials—they have a history! Mineralized eralized skeletons of the Cambrian ancestors were remarkably biological materials are preserved as fossils far more frequently similarto their living counterparts.12 For example, the skeletons than their unmineralized counterparts.In fact, until the early of the Echinodermata have very characteristic sponge-like 1950s it was thought that the fossil record began with the structures made out of relatively large single crystals of calcite. advent of the Cambrian era (about 550 million years ago) when mineralized exoskeletons evolved. We now know that The earliest Cambrian fossils of this phylum have the same J.Mater. Chem., 1997, 7(5), 689–702 689subset of these are common components of endo- and exo-skeletons.6,14 These include the two polymorphs of calcium carbonate, calcite and aragonite, and amorphous silica. Less commonly used minerals are vaterite, amorphous calcium carbonate and phosphate, and crystalline carbonated apatite.The latter is used almost exclusively by the vertebrates and a few invertebrate groups. Quantitatively they do not constitute a significant portion of the biogenic minerals formed today.6,14 Among the biological materials that are formed under relatively controlled conditions, we dierentiate three types based on the organization of their mineral constituents.The one type is composed of multicrystalline arrays, in which the individual crystals are generally all aligned at least in one direction, and often in all three directions. The best known examples of such materials are bones, teeth and shells of various types. In the second type, a single crystal or a limited array of relatively large crystals constitutes the entire structure.The echinoderms are best known for forming such large single crystal skeletal structures of calcite. Many spicules that are used to stien organic structures15 are also composed of single crystals, and again usually calcite. The third group produce biological materials containing an amorphous mineral, the most common being amorphous silica. These structures can vary enormously in size and particularly in shape.Fig. 2 shows an example of each type. The components that perhaps most distinguish biological materials from synthetic materials are the biological macromolecules. They form an intimate mix or composite with the mineral phase at all dierent hierarchical levels, starting at the scale of nanometres. In all the mineralized tissues in which the macromolecules have been even partially characterized, they are found to be very diverse and heterogeneous.In fact, initially it was thought that the macromolecular constituents were more or less unique to each mineralized material. This impression changed substantially more than a decade ago when it was recognized that many of these macromolecules have common chemical attributes—they are rich in carboxylate groups.10 These may be constituents of the protein moieties and/or the polysaccharide moieties.Many of this class of macromolecules also have, in addition to the carboxylate groups, phosphate and/or sulfate groups. The presence of all Fig. 1 (a) Shell of the primitive mollusc Neopilina hyalina showing the these charged groups makes these macromolecules excellent prismatic structure of the adult shell.(b) High magnification view of candidates for interacting with the mineral ions in solution or the aragonitic prisms. Scale bar: 10 mm. with the surfaces of the solid phase.16–18 For convenience, we will refer to the members of this class as ‘control’ basic structures. Similarly, some of the earliest known mollusc macromolecules. shells were composed of prisms;12 a structural motif still Our studies of the control macromolecules of calcium car- common to this phylum today (Fig. 1).13 Even the best pre- bonate-containing biological materials from several phyla sug- served fossils are altered during the passage of time, and we gest that this class can be somewhat arbitrarily subdivided cannot be sure that macroscopic morphological conservation into several groups.One is the aspartic acid-rich proteins and is also indicative of microscopic and molecular structural glycoproteins, which tend to be associated with the crystalline preservation. mineral phases. A second group is the glutamic acid (and/or Studies of modern mineralized tissues show that organisms glutamine)- and serine-rich glycoproteins, which are the major use many dierent minerals and macromolecules, and these components of the several amorphous CaCO3-containing min- are organized into innumerable structural motifs.Bearing in eralized tissues we have examined recently from two widely mind that many of the minerals are actually rather poor diverging taxa.19 A third group is characterized by being building materials, we can guess that with the constant compe- relatively rich in polysaccharides, with proteins containing tition to survive, biological materials were continuously put to fairly average (ca. 10 mol%) amounts of Glx, Asx and Ser. the test and modified to meet the challenges; hence their These macromolecules are the major components of echino- enormous diversity. Identifying the ‘solutions’ they found to derm skeletons and are minor components of mollusc shells.20 these challenges makes their study so special—a theme we will The control macromolecules are usually the quantitatively pursue in this review.minor macromolecular components of a biological material. The major components are more hydrophobic, often cross- Components of mineralized biological materials linked and are hence relatively insoluble in mild acids or at neutral pH.They can vary considerably from tissue to tissue Minerals, macromolecules and water are the major components of these materials. The vast majority of biological materials and in many cases they are indeed tissue specific.6 Unlike the control macromolecules which are dicult to extract or contain only one mineral type.Where two or more minerals are present, they are usually in dierent locations, such as the degrade without dissolving the mineral, these macromolecules can often be extracted or degraded chemically in the presence inner and outer layers of mollusc shells. More than 60 dierent minerals are known to be formed biologically, but only a small of the mineral, implying that they are less intimately associated 690 J.Mater. Chem., 1997, 7(5), 689–702weaker bone with new stronger bone.22 This type of control, like all control, is exercised through the cells directly associated with the tissue. These cells communicate with other cells in order to orchestrate the complex processes of tissue formation. So in a very real sense, the study of the design features of biological materials in general, reveals the ‘intelligence’ and often amazing capabilities of living cells.14 It is therefore unlikely that we can mimic these processes synthetically. We can, however, try to elucidate the design principles and use them to improve our synthetic materials.We will now examine control processes, starting at the A° ngstrom level, and progress through to the millimetre level.Control at the level of A ° ngstroms: from the dissolved molecules to the crystal All mineralization first involves A° ngstrom level processes that start with a solid phase forming from solution. The test case that we choose to follow in detail is calcium carbonate precipitation, either in the crystalline form or as an amorphous hydrated phase.Calcium carbonate crystallizes in five dierent polymorphs, and in addition, an amorphous form. Calcite, aragonite and vaterite are stable under appropriate conditions, while calcium carbonate mono- and hexa-hydrate and amorphous calcium carbonate are very unstable and hence rare in non-biological environments.23 In the biological world, there are very few examples of the monohydrated form,6,24 and no known example of the hexahydrated form.Vaterite is present in some ascidian spicules,25 in a variety of gravity receptors6 and in the egg shells of some gastropods,26 but is, as a whole, also quite rare in biomineralization. The highly unstable amorphous calcium carbonate27 is produced and stabilized biologically, and in fact may be much more abundant in biomineralization than is currently believed.By far the most abundant forms of calcium carbonate produced biologically are calcite and aragonite. It is thus appropriate to consider these two structures in some detail. Calcite and aragonite crystal structures Aragonite and calcite have very similar crystal structures and thermodynamic stabilities.23 The former is slightly less stable than the latter at ambient temperatures and pressures, but is Fig. 2 (a) Tooth enamel of the incisor of a rat. Each elongated rod is very common in biomineralization.6 Both calcite and aragonite composed of hundreds of spaghetti-shaped crystals of carbonated crystal structures are composed of alternating layers of calcium apatite. Scale bar: 10 mm. (b) A ventral plate from the arm of the brittle ions and carbonate ions perpendicular to the c axis (in the ab star Ophiocoma wendti (Echinodermata).The whole structure is one plane) (Fig. 3).23 The calcium ions occupy almost the same single crystal of calcite. Note also the spongy stereom structure changes lattice positions in this plane, and in both structures the in texture in dierent parts. (c) Amorphous silica deposit in the cell walls of the wheat plant T riticum aestivum.Scale bar: 10 mm. carbonate ions lie with their molecular planes parallel to the ab layer. In aragonite, however, some of the carbonate ions are raised in the c direction to form two layers separated by with the mineral phase. They have been referred to as ‘frame- 0.96 A° , and their orientations in the two layers are dierent. work macromolecules’, a term which alludes to their major conceived function, namely providing a three-dimensional matrix in which the mineral phase forms, and a substrate from which some of the control proteins interact with the mineral phase.10 Common examples of framework macromolecules are Gly- and Ala-rich proteins (structurally similar to silk-fibroin) in mollusc shells, type I collagen in bone and tooth dentin, amelogenin in tooth enamel, a-chitin in crustaceans and b-chitin in mollusc shells.6 There are some interesting cases, such as tooth enamel, where the framework proteins are broken down enzymatically and removed during mineral formation.21 This is presumably to allow the crystals to grow larger and form a very dense and mechanically resistant outer layer for vertebrate teeth.In vertebrate bone, certain mollusc shells and echinoderm skeletons, the originally formed mineralized composite material may be locally removed to remodel the material as growth alters its functional requirements, or to replace, in the case of Fig. 3 Crystal structures of (a) aragonite and (b) calcite. Note that the c axis has been tilted out of plane by 5° to improve perspective.bone, older more mineralized and probably mechanically J. Mater. Chem., 1997, 7(5), 689–702 691This shift is the basis for the very dierent properties of these Calcite–aragonite nucleation two phases. An easily conceived way of inducing nucleation of an ionic The optimization of the interactions in aragonite allows crystal is from a cationic plane.This only requires initial better packing, and consequently the density of this phase is concentration and complexation of ions from solution onto a higher than that of calcite. In aragonite growth is preferred matrix substrate with negative charge. Such substrates are very along the c axis, relative to the other crystallographic direc- abundant in biology.This mechanism was shown to operate tions. Thus under conditions of normal temperature and in artificial systems, when crystallization was induced from pressure, aragonite forms as thin needles (acicular crystals) various monolayers of long chain fatty acids deposited at the that do not generally grow into large crystals. Even when they air/water interface. The first reported example involved the appear to do so in some biogenic crystals, the large crystals oriented nucleation of sodium chloride crystals from the homo- are in reality highly twinned, i.e.formed of polycrystalline charged (111) plane under monolayers of stearic acid.35 Mann domains.28–30 Synthetic calcite, in contrast, grows as almost and co-workers36,37 subsequently performed analogous experi- isotropic rhombohedra delimited by a set of equivalent oblique ments on supersaturated calcium carbonate solution sub- faces; {10.4}† in the hexagonal notation.The stability of these phases, whereby oriented crystallization of calcite and vaterite faces is easily understood from the closely packed arrangement from homocharged cation layers was obtained. The oriented of calcium and carbonate ions along the layer.A layer of high nucleation of calcite from polystyrene surfaces decorated with stability (large layer energy that holds the ions together within sulfonate and carboxylate moieties was studied in our group the layer) is, however, always accompanied by a proportionally as a model for the nucleation process occurring in mollusc low attachment energy (the energy that holds parallel layers shell formation.38 A similar mechanism was also shown to together).We note that the sum of the layer energy and operate when acidic proteins extracted from the mollusc shells attachment energy is constant, because it corresponds to the themselves were adsorbed on rigid plastic substrates.39 bulk energy of the crystal.31 Thus, the stability of the {10.4} Nucleation only by concentration of charge should be, at layers in calcite is also the reason for its mechanical weakness, first approximation, non-specific.No repulsion is created and hence the cause of extreme brittleness. The calcite crystal between the cationic crystal layer and the anionic matrix layer, cleaves easily along its {10.4} planes, called ‘cleavage rhombo- even if the positions of the ions in the two layers do not match hedron’ planes, where a crack can propagate along a minimum- perfectly.Calcite has two homocharged calcium planes, (001) energy pathway (with minimum dispersion of energy).32 In and {01.2}, whereas aragonite has one, (001). Furthermore, in contrast, there is no such plane of easy cleavage in aragonite.33 calcite and aragonite the calcium ion positions on the (001) In biology the two polymorphs are used widely as building plane are, as noted, practically identical.If the only driving materials and the choice of polymorph used is almost always force for nucleation was the recruitment of positive ions on a under strict genetic control. It would appear, therefore, that negatively charged surface, the most stable polymorph should one polymorph oers some advantages over the other, even always be formed.Indeed, only calcite was nucleated from the though both have very similar lattice energies and the same (001) plane on acidic macromolecules adsorbed on plastic, composition. Aragonite has the advantage of not having irrespective of whether the nucleating macromolecules had cleavage planes, but has the disadvantage of its small size and been extracted from calcitic or aragonitic mollusc shell layers.39 needle-like morphology.It also has a strong tendency to form It is therefore dicult to conceive that only nucleation of this spherulitic clusters of crystals with high porosity. Calcite, on type can be responsible for polymorph control.the other hand, tends to form larger crystals, but these are Mollusc shells are among the best studied CaCO3-containing biological materials. They are composed of either calcite or very brittle. An examination of the distribution of aragonite aragonite. In some cases both polymorphs are present, but are and calcite among mineralized biological materials does not always separated in dierent layers (Fig. 4).13 Both calcite and produce any simple or clear-cut answers as to the reason for aragonite crystallize from their (001) planes. The same organ- polymorph selection by organisms. ism always produces the same polymorph at the same site. It does appear to be true that when large single crystals of One conceivable strategy could be the involvement of an calcium carbonate are used as skeletal parts, such as in inhibitor of the stable polymorph in solution, while the echinoderm spines and tests and in sponge spicules, they are normally composed of calcite. The large prisms of the prismatic layer of mollusc shells are also usually built out of calcite.Some molluscs do, however, produce aragonitic prismatic layers. There is no obvious advantage or reason for this choice.In contrast, the molluscan nacreous tablets are always composed of aragonite, although very similar structures are produced by some bryozoans out of calcite.34 We know that organisms are able to circumvent the problems arising from calcite brittleness (see the section on control at the nanometre level). It is, however, not at all clear whether organisms ‘relate’ to the calcite–aragonite dichotomy with the same simplistic mechanical analysis as we would deduce from their basic properties. Whatever the reason behind the choice of one polymorph rather than the other, the key step in polymorph determination must be crystal nucleation. Polymorph control during nucleation is thus the next subject to be considered. † The notation {h,k,l} indicates the family of symmetry-related faces or planes. (h,k,l ) indicates only one member of the family.[h,k,l ] Fig. 4 Fracture surface through the shell of the bivalve mollusc,Mytilus indicates the direction of the vector perpendicular to the plane. When the notation (hk.l) is used e.g. (10.4), the period is in the plane of the californianus, showing the outer calcitic prismatic layer (top) and the inner aragonitic nacreous layer (bottom). Scale bar: 10 mm.fourth index, i, in the hexagonal system. 692 J. Mater. Chem., 1997, 7(5), 689–702substrate proteins are responsible only for nucleation and some features of composite materials at the nanometre scale (Fig. 6). This is the scale at which the chitin fibres are orientation of the crystal. The obvious candidate in biological mineralization was considered to be magnesium, which is intergrown intimately with the crystallites.Interestingly, in the artificial assembly the single crystalline domains within the known to favour aragonite formation by inhibiting calcite growth.40–42 In fact aragonite precipitates out of evaporating polycrystalline spherulites of calcite preserve the size and morphology typical of the mineral.Calcite crystallites range sea water because of its high concentration of magnesium. Whatever the controlling element, chemical or structural, it up to 500 nm in size and develop well defined {100.4} cleavage rhombohedron morphologies, while the aragonite crystallites must be present selectively in the microenvironment where the crystal forms.The microenvironment of nucleation is thus the achieve a maximum size of 150 nm and have ill-defined elliptical shapes. key to understanding the process. Mollusc shell nacre is the best studied tissue in this respect. Belcher et al. studied a dierent in vitro system, using as a nucleating matrix the so-called ‘green layer’ sheet isolated from Crystallization occurs inside a pre-deposited matrix,43 composed of thin layers of b-chitin sandwiched between two thicker abalone shells.49 They also observed aragonite crystallization when proteins extracted from abalone shell aragonitic phase layers of silk fibroin-like proteins, onto which acidic macromolecules are adsorbed.44,45 The fibre axis of the chitin and were added to the green layer, and calcite crystallization when calcite-extracted proteins were added.silk proteins are perpendicular to each other, and aligned with the a and b axes of the aragonite tablets, respectively Nucleation of calcite and aragonite from the (001) plane is common in biomineralization. Well studied examples are, in (Fig. 5).46,47 This well defined spatial relation between substrate and overgrowth phase suggests an epitaxial mechanism of addition to mollusc shells, coralline algae,50 calcareous sponge spicules51 and sea urchin larval spicules.52 There is also evi- nucleation. Surprisingly, the same mollusc shell acidic macromolecules dence of oriented nucleation of calcite from the homocharged (01.2) layer.This occurs in certain scimitar-shaped calcareous that exclusively induced calcite formation when adsorbed on plastic substrates, were shown to retain polymorph specificity sponge spicules.53 The calcite crystal forming the spicule is oriented such that the [01.2] direction is always along the in an appropriately assembled artificial microenvironment, designed to match roughly the biological one.48 The acidic spicule axis.In addition, the c axis direction is uniquely fixed such that the positive end always points out of the convex glycoproteins associated with calcitic prismatic and aragonitic nacreous layers of various mollusc shells were adsorbed on an part of the spicule.The combination of the asymmetric spicule morphology and its uniquely defined relationship to the crystal artificial assembly of b-chitin (from squid pen) and silk (from silkworm cocoons). Neither of these matrix components is axes orientation can only be explained if the nucleation surface structure is totally controlled.This includes distinguishing calcified in the original tissue. Once adsorbed on this scaold, the macromolecules extracted from aragonitic mollusc shell layers induced aragonite formation, while those extracted from calcitic layers induced calcite formation, with total fidelity.When no acidic macromolecules were introduced, only vaterite spherulites formed on the chitin surface layers. The orientation of the nucleated crystals relative to the inducing proteins is not yet known. If these crystals are nucleated from the (001) plane, similar to their orientation in vivo and in vitro after adsorption on plastic, it would be tempting to conclude that a three-dimensional nucleation site fixes the carbonate positions, in addition to those of the calcium ions.The structural requirements for such a nucleation site, however, appear to be prohibitively stringent. Another possibility is a combined nucleation–inhibition mechanism, but in this case the inhibition and nucleation must involve one or more proteins.It certainly does not involve magnesium which was absent in the experiment. At present we do not understand the mechanisms involved in vitro, and certainly not in vivo. The resulting calcite- or aragonite-impregnated chitin, although not as well organized as in mollusc shells, possesses Fig. 5 Schematic block diagram showing the spatial relations between Fig. 6 Synthetic composite materials produced in vitro containing the crystallographic axes of an aragonitic nacreous polygon and the underlying organic matrix. [From On Biomineralization by H. A. (a) calcite and (b) aragonite crystals in a matrix.The matrix is composed of b-chitin and silk fibroin, as well as soluble proteins from the calcitic Lowenstam and S. Weiner.Copyright © 1989 by Heinz Lowenstam and Stephen Weiner. Used by permission of Oxford University Press, shell of Atrina serrata in the case of (a) and from the aragonitic shell of Elliptio sp. in the case of (b). Inc. (ref. 6 of this work)]. J. Mater. Chem., 1997, 7(5), 689–702 693between the positive and negative surfaces of the (01.2) layer. imposed directly or indirectly by surrounding specialized cells.In the simplest scenario, crystals grow in a solution containing This is equivalent to saying that the nucleated surface is chiral and defined in three dimensions. Note that the calcite structure the component ions, not dissimilar from conventional growth from solution in non-biological environments. The crystals is not chiral. Mann and Sparks pointed out an analogous case in coccoliths,54 where the morphology of the single crystals is tend thus to assume their regular growth morphology.Growth is stopped only by contact with neighbouring crystals. This is asymmetric and uniquely defined, suggesting chiral recognition at the nucleation stage. Furthermore, Berman et al.55 recently probably the case for calcite crystals in egg shells59 and aragonitic crystals in scleractinian corals,60 and fish otoliths6 studied the nucleation of calcite under monolayers of polydiacetylene carboxylates.They observed nucleation from the that grow as polycrystalline bundles o predisposed nucleation centres (Fig. 7). As orientation is not well controlled during (01.2) plane, as well as orientation within the plane relative to the polymer backbone direction.This implies that almost nucleation, a less tightly packed material is formed, which is porous and brittle. The bulk material produced reflects this complete control over the nucleation site geometry may be achieved under artificial conditions. Whether there is any process, and is thus relatively weak.1 The properties of the material are therefore controlled to a large extent at the level advantage in such a high level of nucleation control in biology is not clear.What is clear is the enormous intrinsic controlling of nucleation, by the density and the relative geometry of the nucleation sites. In mollusc shell nacre and simple prismatic power of some of the biological nucleation processes. layers only one well oriented crystal originates from each nucleation site.The lateral growth of both aragonite and calcite crystals is also limited only by meeting the neighbouring Control at the nanometre level: crystal growth and growing crystals, resulting in a typical honeycomb structure of morphology irregular polygons (Fig. 8). There are many examples in biomineralization where single The next step, following crystal nucleation, is the growth of the crystals into desired shapes and sizes.Crystals grow by crystals grow as separate entities with well defined individual morphologies and sizes that are very dierent from their non- progressive addition of molecules or ions onto the crystallization nucleus. Growth in the various directions is governed biological counterparts.All these crystals grow inside closed kinetically by rules determined by the crystal structure and symmetry. Molecules will be added faster where the balance of the interactions with the existing crystal is more favourable. In general, adding a molecule within a growing crystal layer is more favourable than creating a new layer. The first molecule of a new layer makes contacts only with molecules of the underlying layer, while a molecule added at a growing step or kink establishes contacts in two or three directions. Thus crystals normally grow in layers, and are delimited by a well defined set of faces.Spherical smooth surfaces are only observed above the so-called roughening transition, where the driving force to growth is so large that adding a molecule in any position does not make, kinetically, any dierence.56 The growth morphology of crystals is determined by their relative rates of growth in the various directions.For example, if it is much easier to add molecules in one unique crystallographic direction relative to all others, the crystal will develop as a needle. On the other hand, if the energetics involved in adding layers of molecules in all directions are approximately equivalent, the resulting crystal will be roughly isotropic in shape.The slow growing directions are the ones that determine the crystal morphology, with the layers perpendicular to them developing as stable faces (having high layer energy and low attachment energy).31 Each crystal thus has a typical growth morphology under a given set of conditions.These include physical parameters such as temperature, pressure and supersaturation, and chemical parameters, such as interactions with the solvent and with cosolutes. In particular, both co-solutes and solvent may act as inhibitors of crystal growth in specific directions.57 If they are adsorbed on certain crystal planes rather than others, crystal growth will be slowed down in the directions perpendicular to the planes.A set of faces parallel to the plane may consequently develop, or increase in morphological importance, when already present. Macromolecular inhibitors, that structurally match the molecular motif on one set of crystal planes, may interact with these planes from solution in a manner equivalent to the process described for nucleation. This results in modulation of crystal morphology through the above mechanism.18 In biology, the microenvironment where crystallization occurs is the key to the control over crystal growth, as well as nucleation.Crystals are generally formed in pre-defined spaces, delimited by extracellular matrices and cell membranes, or Fig. 7 Fracture surfaces of (a) the calcitic egg shell of the domestic inside vesicles.14,58 Inside these defined spaces the crystals grow hen, and (b) an aragonitic otolith from the bony fish Seriphus politus (reproduced with permission from ref. 18, p. 158). Scale bars: 0.1 mm. under shape, size, concentration and composition constraints 694 J. Mater. Chem., 1997, 7(5), 689–702measured the Young’s modulus of sea urchin larval spicules and indeed showed that it is quite dierent from pure calcite.63 We, however, suspect that this may be also due, in this particular case, to the presence of some amorphous calcium carbonate79 (see stabilization of amorphous calcium carbonate).Sea urchin spine It has been long recognized that sea urchin spines are each composed of one single crystal of calcite (based on polarized light and X-ray diraction), with the c axis of the single crystal oriented along the morphological axis of the spine.64 The single crystal grows inside a membrane (syncitium) in communication with many cells that provide the ions and all other biological components necessary for crystal growth and shaping.65,66 The Fig. 8 Fracture surface of the prismatic calcite layer of the shell of the cells populate the meanders of the channels (stereom) running mollusc Atrina serrata showing the polygonal crystals.Scale bar:10 mm. all along the spine in a continuous structure. The spine grows by elongation at the tip and thickens on all the peripheral surfaces. The result is a convoluted spongy element. This is spaces delimited by lipid bilayer membranes or macromolecu- later filled in with mineral, giving rise to a radial structure of lar matrices.It could be envisaged that shape is established full sectors of calcite, connected by spongy septa (Fig. 9). The directly and only by the membrane or matrix, by simple mature spine still diracts X-rays as a single crystal. The mechanical interference. It is, however, necessary to invoke a presence of channels not only provides a means for the cells mechanism whereby an intrinsically soft and deformable bar- to populate the whole spine, but also contributes to the rier can overcome the forces acting on it by the growing mechanical performance of the material. Spongy structures are crystal.Other mechanisms are conceivable, that probably both lightweight and more elastic than full structures.1 operate together with the membrane/matrix to achieve final However, the typical size range of a septum in the stereom, shape determination.These are induction of growth in con- ca. 1 mm, is still very large relative to the size of the unit cell trolled directions and active growth inhibition. In the induction of the crystal. Fracture of the septa could thus still easily occur scenario, the component ions may be delivered into the crys- along the cleavage planes of calcite, but in fact does not tallization space at specifically controlled sites, such that (Fig. 10). The organism adopts the reversed composite material growth can occur only in certain directions. An example is the approach to further reinforce the crystal against fracture.67,68 sea urchin larva where calcium preferentially enters the vesicle Glycoproteins are trapped inside the spines in amounts of close to the fast growing tips of the spicule.61 This would imply ca. 0.02% by mass of mineral. New calcite crystals, grown a close proximity of the ion pumps presumably in a membrane epitaxially on the cleaned spines develop, in addition to the with the growing mineral, at least during the final stages of stable {10.4} faces, a set of unstable faces, slightly inclined to growth. It has also been observed in sea urchin larval spicules the c axis [of index {01,l }, with l#1.5].69 The original spine that the just-nucleated crystals display the regular {10.4} faces of calcite.62 The formed spicules, however, always terminate with smooth and curved surfaces.The formation of such curved surfaces is in itself dicult to understand. It probably requires some other mechanism that keeps all the growth sites on the crystal surface active, similar to the situation occurring above the roughening transition.56 One possibility, but by no means the only one, would be an inhibition process that, by interfering continuously with the completion of the crystal layers, would generate a surface composed of steps in the nanometre scale.Such inhibitors may even be active during the entire crystal growth process. In the inhibition scenario for controlling shape, growth inhibitors are delivered into the solution and are adsorbed actively onto the growing crystals in controlled directions. This mechanism presents an additional attractive possibility, of actively modifying the properties of the crystal bulk, while modulating its shape.Some of the inhibitors adsorbed at the crystal surface are eventually overgrown, and remain occluded inside the crystal, specifically along the planes where they have been adsorbed. If these are sheet-structured macromolecules, and adsorption occurs with a high enough frequency, the final result is a kind of fibre-reinforced reversed composite material. 18 The host crystal constitutes a continuous matrix that is hard and often brittle. The guest macromolecules embedded inside it are the fibres or sheets that endow the crystal with pliancy and increased resistance to brittle fracture. In the case of calcite one may envisage that any type of interference with the propagation of cracks along the cleavage planes would reinforce the crystal against fracture, by both deviating and absorbing the propagating crack energy.32 This mechanism Fig. 9 Fracture surfaces of (a) immature and (b) mature spines of the appears to be exploited as a reinforcement strategy by organ- sea urchin Paracentrotus lividus. Illustration (a) is reproduced with permission from ref. 18, p. 159. Scale bars: 0.1 mm. isms that choose to build single crystal skeletal elements. Emlet J. Mater. Chem., 1997, 7(5), 689–702 695that separate contiguous imperfections (the coherence lengths). They aect maximally the width of the diraction peaks from a set of planes containing the imperfections, but not from planes forming a wide angle with them.Three-dimensional mapping of the distribution of imperfections in ten sets of biogenic calcite single crystals of very dierent shapes (sea urchin spines and larval spicules, five dierent kinds of calcareous sponge spicules, single prisms from mollusc shells and two kinds of foraminifera shells) showed in seven cases out of nine a striking correspondence with macroscopic crystal shape.53,70,72 There is thus a link between textural properties at the nanometre level and crystal shape at the sub-millimetre level.One possibility we proposed is that the macromolecules shape the growing crystal by specific adsorption onto some crystal faces and not others. One exception is the so-called slender monaxon spicule from the calcareous sponge Sycon, that does not contain any occluded protein.Its texture is isotropic, as is the texture of pure calcite. The second exception is the prisms from the shell of the mollusc Atrina. The prisms are elongated along the c axis, but the coherence length is shorter along c, indicating higher protein intercalation in that direction. This is, however, also the only case we studied of a single crystal taken from a polycrystalline assembly (the prismatic layer), where crystal growth occurs in a preformed organic matrix.Growth in the lateral directions is stopped by the matrix and/or by the adjacent crystals. Furthermore, we have independent proof that the main components of the intracrystalline macromolecules, proteins rich in aspartate, are indeed intercalated along the (001) planes.We thus conclude that in Atrina a dierent mechanism is operating in the determination of crystal Fig. 10 High magnification views of the fracture surfaces of (a) the morphology. calcitic sea urchin spine showing conchoidal cleavage, and (b) a The correspondence between coherence length distribution synthetic calcitic crystal showing the smooth surfaces of the {10.4} and shape is particularly striking for the curved monaxon cleavage planes.Scale bars: 10 mm. spicules and asymmetric triradiate spicules from the calcareous sponge Sycon.72 As noted, the scimitar-shaped curved monaxon concomitantly becomes etched, suggesting that glycoproteins is elongated in the general direction [01.2]. The circular leaked out of the etched surface and readsorbed at the growing section of the spicule contains many non-equivalent crystallo- crystal surfaces along the {01, l} planes. In agreement with graphic directions, and correspondingly the coherence lengths this interpretation, calcite crystals grown de novo from a (lc) are almost identical, lc#1500 A° .On the other hand, of the solution containing the same glycoproteins released from the three equivalent {01.2} reflections, the one along the morpho- spines after dissolution, developed the same morphology as logical axis of the spicule has lc#8000 A°.The other two, the overgrown crystals. The crystals grown in the presence of inclined to the morphological axis by 60°, have lc#2000 A° the glycoproteins are also mechanically more resistant to (Fig. 11). This phenomenon can only be explained by assuming fracture than pure calcite. They cleave with a conchoidal- an accurate, nanometre-scale controlled delivery of the proteins type fracture similar to the biogenic spines, and are very dierent from the pure calcite crystals.67 The latter shatter easily under an applied force, with the fracture lines always being along the cleavage planes of calcite (Fig. 10). Recently, Albeck et al.20 showed that the key constituents of the glycoproteins that interact with the growing crystals involve oligosaccharide chains linked to the polypeptide chain in tightly structured clusters. Modulation of crystal texture To gain more insight into how macromolecules are occluded inside single crystals, we have mapped by X-ray diraction the microtextures of a series of biogenic single crystals of calcite from various organisms.70 A macromolecule is far too large to be incorporated into the perfect lattice of a single calcite crystal.Thus, when it is adsorbed and eventually overgrown by the growing crystal, its presence will leave a permanent imprint inside the crystal, in the form of an imperfection. Imperfections always exist even in the most perfect crystals.71 Their distribution can be characterized by means of the diraction behaviour of the crystal.Diraction originates from domains of perfect structure, and the sharpness of the dirac- Fig. 11 Two views of the calcitic curved monaxon spicules from the tion peaks is inversely proportional to the size of the perfect calcareous sponge Sycon sp.The lengths of the superimposed arrows domains. Imperfections intercalated along certain crystallo- are proportional to the coherence lengths in the crystallographic directions indicated. graphic planes limit the size of the domains to the distances 696 J. Mater. Chem., 1997, 7(5), 689–702onto the growing crystals. Interestingly, protein intercalation under ambient conditions, and therefore minimal energy is required for this aspect of its formation. This contrasts with is mirrored by the mechanical properties.Microindentation performed on polished longitudinal sections of the spicules the other fairly commonly used amorphous minerals, which do need to be stabilized. Amorphous calcium phosphate is results in anisotropic crack propagation along the spicule, in the same unique direction where proteins are not intercalated.73 often used for temporary storage of ions, because its solubility is higher than that of the crystalline materials.It is also used, It would thus appear that protein intercalation serves the double purpose of modulating shape and mechanical proper- however, for skeletal strengthening purposes, for example in some ascidian spicules and the gizzard plates of some ties.When the crystal morphology matches the crystal symmetry, it may be sucient to exploit the recognition capabilities gastropods.6 Amorphous hydrous iron(III ) phosphate is the mineral used in sternal shields of certain annelids6 (Fig. 13). of the (glyco)proteins for specific crystal motifs. When, however, the single crystal morphology does not respect the crystal Amorphous calcium carbonate is also formed by several organisms in widely divergent taxa.6 It is most abundant in some symmetry, the targeting strategy can overcome the intrinsic anisotropy of the crystal, and of the protein–crystal interactions plants, where it presumably functions as a temporary storage site for ions.77 It is used for structural purposes, such as in the as well.This raises the intriguing question of whether or not this biological ‘override’ of the inherent nature of the crystal– spicules of ascidians of some Pyuridae78 (Fig. 14), in the spicules of the sponge Clathrina,19 and as a precursor phase of protein interactions has the functional purpose of producing a more isotropic material in terms of defect distributions.calcite in sea urchin larval spicules.79 We elaborate briefly on the case of amorphous calcium carbonate, not because it is so Another interesting illustration of this strategy was observed recently in sea urchin spines. The diraction data indicate that abundant in the field of biomineralization, but because it presents such intriguing paradoxes.the anisotropy in crystal texture (c vs. ab) is larger in mature secondary spines, where the stereom is already filled with The use of amorphous calcium carbonate is puzzling. The mineral is very unstable, and its transformation into one of mineral sectors, than in immature spines that had only developed the spongy stereom. Etching of broken stereom sections the crystalline polymorphs is extremely fast in solution under normal conditions.27 Organisms must invest a lot of energy to of immature spines show curved, onion-like mineral deposition lines, transverse to the septa, irrespective of their direction stabilize this phase, and hence presumably derive considerable benefit from using this unusual mineral.The strategy used for (Fig. 12).74 These lines do not appear in the filled sectors, suggesting a dierent, possibly less controlled, mechanism of stabilization of amorphous calcium carbonate again involves specialized macromolecules.Recently, glycoproteins have been crystal growth during the filling stage. Interestingly, synthetic calcite crystals grown from solution in the presence of the isolated from within the amorphous mineral of both Pyura antler spicules and Clathrina triradiate spicules. Their amino proteins extracted from the spines have even higher textural anisotropy. This is true both for growth along c relative to the acid compositions, rich in glycine, serine and glutamic acid, are very similar.They both have associated oligosaccharides. ab plane, and, within the ab plane, between the directions [10.0], where protein intercalation occurs, and [11.0].The When introduced into supersaturated solutions of calcium carbonate, both prevent crystallization completely, and the mechanism of growth during the filling stage in the spine is thus closer to that of the protein-containing synthetic crystals, where no control over the microenvironment is exercised.75 The ‘strive for isotropy’ may thus be a more widespread strategy in the construction of single crystal skeletal elements.Stabilization of amorphous calcium carbonate If the achievement of isotropy in mechanical performance is an important issue, the best construction material should in itself be intrinsically isotropic. This property is shared by amorphous minerals. Amorphous silica is indeed used by a wide range of organisms, from the complex beautifully sculpted diatoms to siliceous sponge spicules and plant phytoliths.76 In terms of quantities formed worldwide, silica is one of the three most abundant biogenic minerals (together with calcite and aragonite).It therefore appears to oer important benefits as a component of biological materials. In addition to being isotropic, silica has the obvious advantage of being stable Fig. 13 Sternal shields of the marine annelid Sternaspis sp. composed of an amorphous hydrous iron(III ) phosphate. Scale bar: 0.5 mm. Fig. 12 EDTA-etched fractured surface of an immature sea urchin Fig. 14 Antler-shaped spicules of the marine ascidian Pyura pachydermatina composed of amorphous calcium carbonate spine showing the mineral deposition lines J.Mater. Chem., 1997, 7(5), 689–702 697amorphous precipitate that is consequently formed is stable over long periods of time.19 Control at the micron level: the intimate involvement of cells Cells form biological materials. Their involvement can be direct, with the mineralization structures forming in specialized vesicles within the cells, or in close association with cell walls.It can also be indirect in that the cells synthesize and release Fig. 15 X-Ray diraction patterns of the aragonitic nacreous layer of macromolecules to the extracellular environment. Here they the bivalve Neotrigonia margaratifera. Hundreds of crystals are in the self-assemble into a three-dimensional framework or matrix in diracting volume. The patterns in two orthogonal directions show which the mineral subsequently forms.Whatever the process that they are all relatively well oriented in three directions. used, one level of structural organization of a mineralized biological material can frequently be related to the size of the hundreds of microns have shown that in gastropods there is cells that form the structure.In general such cells tend to be no lateral preferred orientation whatsoever. In the shelled elongated and range in size from a few microns to ten microns cephalopod, Nautilus, there is some degree of preferred orien- in cross-section and can be tens of microns long. It is the tation, as is also the case in many bivalves.86,87 For example, cross-sectional plane which usually interfaces with the extra- in the bivalve Neotrigonia margaratifera the extent of organiz- cellular environment.This length scale may constitute a ‘struc- ation can be rather good in all three directions45,87 (Fig. 15). tural benchmark’ of cellular activity, and is often a key element These observations show that the cell influences the size and in the structural organization of a biological material.orientations of the crystals in two dimensions. The third Cellular controlled organization at sub-micron levels can be dimension (layer thickness) is presumably determined by the imposed by the scale of the spaces in the three-dimensional properties of the self-assembled matrix. These studies also extracellular matrix, or by the formation of the mineralized show that the cells determine the orientations of the crystallo- building blocks in vesicles within a cell, followed by assembly graphic axes indirectly through the matrix substrate.The fact outside the cell. Examples of the latter are the marine cal- that the degree of orientation of whole areas of polygonal careous plants Coccolithophoridae80 and the marine proto- crystals is genetically controlled suggests that there may be zoans belonging to the group of the miliolid foraminifera.81 some selected mechanical advantage for random crystal organ- One phylum which consistently forms single crystals that are ization in one case vs.preferred orientation in others. much larger than the size of normal cells is the Echinodermata. The nacreous layer functions mechanically as a classic Their strategy is to have a whole team of cells fuse their composite material rather than a ceramic, despite the fact that membranes to form a giant vesicle or syncitium.6 A single the organic component usually constitutes only ca. 1% by calcite crystal is nucleated within the syncitium and in some mass of the material. It is also a platelet-reinforced composite, cases can grow to even centimetre size (see Sea urchin spine).as opposed to the more common fibre-reinforced composites Here we will examine the product of cellular activity on the of the synthetic world.88 Mechanical studies demonstrate well higher order structural organizational patterns of two well the rather remarkable bulk materials properties of nacre both studied mineralized materials, the mollusc shell nacreous layer under compression and under tension.89,90 Observations of and bone.fracture planes show clearly the tortuous route followed by the crack as it progresses along the matrix sheets or in the Mollusc shell nacreous layer perpendicular direction as it traverses across the crystal layer between tablets. At this structural level, the nacre is deduced The cells that form the nacreous layer are located on the side to derive its unusual mechanical properties directly from its of the shell-forming tissue (the mantle) that faces the inner highly ordered layered structure, prompting the conclusion surface of the shell.They are usually close-packed and hence that no really novel mechanisms are involved in achieving its polygonal in cross-section.82 These cells form an extracellular mechanical properties.88 We suspect, however, that this may matrix in which the aragonitic crystals grow.The dominant not be the case. We note the unique plywood-like structure of matrix structural feature is a series of sheets regularly spaced the matrix itself, the fact that it is composed of two very at distances of a half to one and a half microns from each dierent polymers (chitin and silk fibroin-like protein), the other (Fig. 4). The resultant mineralized structure is composed very real possibility that macromolecules are also occluded of polygonal-shaped flat tablets of aragonitic crystals, with inside the aragonitic crystals where they may alter the bulk each layer of crystals separated by a matrix sheet.Although it properties of the mineral phase, and the well designed interface has still not been demonstrated directly that each polygonal between matrix and mineral inferred from the documented crystal is formed by one mantle cell, the observed correspon- specific spatial relations between them. All or some of these dence in size between crystals and cells in dierent species features may indeed constitute ‘novel’ design strategies that suggests that this is the case.83 contribute to the unique mechanical properties of nacre.The sheets of matrix formed by the mantle cells are composed In general molluscs oer a wide variety of opportunities to of no less than five dierent layers, following the model investigate structural design features. The commonly formed proposed by Weiner and Traub.45 Each cell probably makes crossed-lamellar structure comprises a three-dimensional array its own three-dimensionally ordered ‘patch’ of matrix and of closely packed aragonitic crystals.The structure is in itself mineral. Atomic force microscope84 and electron diraction fascinating, and the few mechanical studies performed to date studies85 of the vertical orientations of nacreous crystals from point to interesting bulk properties.91 One enigmatic obser- several bivalves show that stacks of four or five layers of vation is that the hardness of the aragonitic shell is greater crystals may be very well oriented. This could be the result of than inorganic aragonite.The matrix component of these shells each stack being nucleated once and a single crystal growing is ca. 0.5% by mass.91 through the matrix sheets. Alternatively, it could be the product of the synchronized activity of a single mantle cell forming a Bone highly ordered matrix–mineral structure. X-Ray diraction studies of the lateral orientations of the a and b crystallographic The basic building block of bone (and tooth dentin) is the mineralized collagen fibril.92 In the world of biomineralization axes of an assemblage of aragonitic tablets extending for 698 J.Mater. Chem., 1997, 7(5), 689–702stage the crystal growth dynamics dominate, and ‘push’ the collagen molecules aside. In mineralized tendons and parallel-fibred bone the extruded fibrils are arranged into long parallel arrays, with the fibril axes all in the same direction.103 In dentin the extruded fibrils are all in the same plane, but are not well oriented with respect to each other.104 The most complex form of bone is lamellar bone.Here the cells extrude the fibrils such that all the fibrils that constitute one newly formed layer are aligned in one direction in a plane. The next fibril layer is rotated by some degree such that a plywood-like structure is formed.The cells control not only fibril orientation, but also the azimuthal orientation of the crystal layers around the fibril axis. These too are rotated with each additional layer. The cells form a complex structured unit 2 to 3 mm thick, and then begin the whole process again.105,106 The resulting so-called ‘rotated plywood’ structure is thus a highly complex composite material (Fig. 17). A detailed study of the microhardness properties of parallel- fibred bone107 by indentation clearly reflected the anisotropic nature of the array of aligned mineralized collagen fibrils. The lowest values are obtained when the indenting direction is perpendicular to the alternating layers of triple helical molecules and crystals (P).It is highest when the crystals are indented edge-on in the direction parallel to the long axis of the bone (T) (Fig. 18). When the microhardness properties of the lamellar bone structure were probed, they revealed the well known general tendency for the bone to be somewhat harder in directions parallel to the bone long axis as compared to directions perpendicular to the long axis. The dierences were, however, gradual when the structure was probed in many dierent directions, and relatively small compared to parallel- fibred bone. It thus appears that the design motif of lamellar Fig. 16 (a) Transmission electron micrograph of an unstained mineralized collagen fibril from calcified turkey leg tendon. Most of the bone is to form a mineralized structure that tends towards plate-shaped crystals of carbonated apatite are viewed face-on.The isotropy, even though the building block used is highly aniso- characteristic 67 nm banding of collagen is also apparent. (b) Schematic tropic. This is achieved by the formation of complex higher illustration of the organization of the crystals in layers in the collagen ordered structures. fibril. Reproduced with permission from ref. 112. Conclusions: towards the millimetre scale and this is a most unusual matrix–mineral composite in that the beyond carbonated apatite crystals are among the smallest, if not the smallest, biologically produced crystals known. They are on Measurements of the mechanical properties of biological mate- the average 50 nm×25 nm×2 nm.6,93 Most of these plate- rials, millimetres in size or larger, can be made relatively easily.shaped crystals are located inside grooves or channels within the type I collagen fibril (Fig. 16) to form a layered structure across the fibril.94 Thus the mineralized collagen fibril is itself crystalline and highly anisotropic. Cells synthesize the collagenpolypeptides and these assemble into small fibrils in vesicles within the cell.These are then packaged for secretion. Further assembly occurs into bundles in the extracellular environment, presumably in such a way that the three-dimensional orientation of the fibrils is well controlled.95 Mineralization takes place in the extracellular environment. In some fast-forming tissues the first crystals form inside very small vesicles.96 These crystals have no preferred orientation.As these mineral-filled vesicles have also been observed in the proximity of the sites of ordered nucleation that occurs within the collagen fibril,97 it is conceivable that they function as a supplier of ions for intrafibrillar mineralization.98 The crystals that form within the fibrils nucleate at a very specific location within the fibril,99,100 and then grow rapidly along their c axes.The latter are well aligned with the collagen fibril axis. At this initial stage the crystals are needle-shaped. They soon, however, grow into plates filling the collagen fibril channels.101 The plate-shaped crystals finally push their way Fig. 17 Fracture surface of lamellar bone from the midshaft of a rat out of the fibril channels into the overlap zone between layers tibia showing several individual lamellar units (top).Schematic illus- of triple-helical molecules.102 Thus the collagen fibril seems to tration of the orientations of the collagen fibrils (cylinders) and the fulfil a matrix framework function by defining the nucleation crystal planes inside them at three dierent locations within a single lamella (bottom).The structure in area 4 is unclear. site location and controlling initial crystal growth. At a later J. Mater. Chem., 1997, 7(5), 689–702 699relatively active mammals, suggesting that it is able to withstand a wide variety of mechanical challenges. The study by Ziv et al.106 of the microstructure–microhardness relations in lamellar and parallel-fibred bone, showed that the former tends to be significantly more isotropic than the latter at the tens of micrometres scale.This observation raised the interesting question of whether other multifunctional biological materials are also structured in such a way as to emphasize isotropic properties. We have noted that this might well be the case for the echinoderm stereom structure at the nanometre scale (see Modulation of crystal texture, earlier). The structure of the shell plates of a most unusual marine barnacle, Ibla, is interesting in this respect.The barnacles are members of the Arthropoda, and generally have mineralized Fig. 18 Schematic illustration of the nanometre-scale structure of the calcitic exoskeletons. Ibla is the exception. It produces a shell mineralized collagen fibril showing the triple helical molecules of plate mineralized with carbonated apatite, the same mineral collagen (cylinders) and the plate-shaped crystals.The arrows show present in bone. The framework constituent of the matrix is the three directions of indentation. (Reproduced by permission of the a-chitin, like all other arthropods. A detailed study of the shell publisher from V.Ziv, H. D. Wagner and S. Weiner, Bone, 1996, 18, 417 (ref. 107 of this work). Copyright 1996 by Elsevier Science Inc.) plate structure revealed remarkable similarities to lamellar bone, right down to the nanometre level109 (Fig. 19). This appears to be an example of convergent evolution producing a very similar, in this case probably more generally functional, Their interpretations in terms of structure, mechanical behaviour and function, however, are dicult, because they incorpor- material in two quite dierent phyla.Isotropy in a material has obvious advantages. ate the contributions of dierent hierarchical structural levels. Many bulk measurements of biological materials have been Macromolecules that constitute the matrix in biological materials are always highly anisotropic, as are the crystalline made and analysed in terms of known mechanical engineering properties.Indeed these are the studies that have shown just mineral components. The substitution of a crystalline mineral how mechanically interesting many biological materials are, especially when compared to analogous synthetic composite or ceramic materials. By analysing the structural properties of biological materials at dierent length scales, it is clear that organisms have evolved a variety of interesting strategies to improve the mechanical properties.This is particularly impressive when bearing in mind the many disadvantageous properties of the starting mineral components. It is particularly helpful to be able to measure directly the mechanical properties at the appropriate length scale of the structural feature of interest, and in particular the key properties that are important for the organism.Unfortunately in many cases, the appropriate tools for making such measurements are not available, and we do not know for sure what the important parameters are. It is also often tacitly assumed or implied that the biological materials are well, or even perfectly, adapted to the needs of the organisms that produce them.This is in reality almost impossible to demonstrate. A dierent conceptual approach to the analysis of structure– mechanical function relations in biological materials millimetres or larger in size, is to dierentiate, if possible, between those materials that are used for many purposes, the ‘concretes’ of the biological world, and those that are structurally designed for specific tasks.The latter tend to have bulk structures and architectures that vary within a given phylum even at relatively low taxonomic levels. A good example is the mollusc shell. Molluscs produce seven major structural materials for their shell layers. These, however, vary from group to group (about 50 variants are known) and dierent structural types are often combined in one shell.108 The overall impression is that each shell type has evolved to fulfil specific functional requirements.A good example of an ‘all-purpose’ type material is the echinoderm stereom structure. The same calcitic material, which has sponge-like microarchitecture, is used by almost all members of this phylum for a wide variety of purposes.6 Another example of a more generally functional material is the chitinous exoskeleton of the arthropods.It has a complex lamellar structure with a well defined plywood motif. In crustaceans it is also mineralized with calcite. This basic skeletal material is used by almost all members of this huge Fig. 19 Fracture surfaces of (a) the shell plateof the marineinvertebrate phylum.6 A third example of such a material, in our opinion, barnacle, Ibla, and (b) lamellar bone from the tibia of a rat. Note the remarkable similarity in lamellar structure. is lamellar bone. It is used by many mammals and in particular 700 J. Mater. Chem., 1997, 7(5), 689–7029 S. M. Stanley, Proc. Natl.Acad. Sci. USA, 1973, 70, 1486. for an amorphous mineral should contribute significantly 10 S. Weiner, W. Traub and H. A. Lowenstam, in Biomineralization towards improving the overall isotropic properties of a and Biological Metal Accumulation, ed. P. Westbroek and E. W. biological material. Biology does indeed make use of a variety de Jong, Reidel, Dordrecht, 1983, p. 205. of amorphous minerals for structural purposes (see 11 S.Bengston and S. Conway Morris, in Origin and Early Evolution Stabilization of amorphous calcium carbonate, earlier). The of theMetazoa, ed. J. W. Schopf and C. Klein, Plenum, New York, 1992, p. 447. full extent of this phenomenon is still probably grossly under- 12 S. Bengtson, S. Conway Morris, B. J. Cooper, P. A. Jell and estimated, as the presence of amorphous minerals is often B.N. Runnegar, Mem. Assoc. Australasian Paleontologists, 1990, dicult to detect, especially when crystalline material is also p. 9. present. 13 O. B. Bøggild, K. Dan. V idensk. Selsk. Skr. Naturvidensk. Math. Silica is the most common and quantitatively most abun- Afd., 1930, 9, 233. dantly formed biogenic amorphous mineral. The sizes of these 14 K.Simkiss and K. M. Wilbur, Biomineralization. Cell Biology and Mineral Deposition, Academic Press, San Diego, 1989. biogenic siliceous products vary from several microns to tens 15 M. A. Koehl, J. Exp. Biol., 1982, 98, 239. of centimetres in the case of some sponges.76 Macromolecules 16 A. Veis, in Biomineralization. Chemical and Biological are associated intimately with these siliceous biological mate- Perspectives, ed.S. Mann, J. Webb and R. J. P. Williams, VCH, rials, both within the mineral phase and as framework struc- Weinheim, 1989, p. 189. tures that order the spherical mineral particles into higher 17 C. S. Sikes, A. Wierzbicki and V. J. Fabry, in Biomineralization order structures.110 The probable reasons why silica is so ’93, ed. D.Allemand and J-P. Cuif, Bull. e’Institut Oce�anog., Monaco, special no. 14, 1, 1994, 1. widely used biologically are that it is very insoluble at neutral 18 L. Addadi and S. Weiner, Angew. Chem., Int. Ed. Engl., 1992, or close to neutral pH, it is relatively abundant in a soluble 31, 153. form in ground water and in sea water (except where diatoms 19 J. Aizenberg, G. Lambert, L.Addadi and S. Weiner, Adv. Mater., have used almost all of it), and it polymerizes rather easily 1996, 8, 222. into a solid phase under a variety of conditions. It is not nearly 20 S. Albeck, S. Weiner and L. Addadi, Chem. Eur. J., 1996, 2, 278. as obvious why the two other fairly common biologically 21 C. R. Hiller, C. Robinson and J. A. Weatherell, Calcif. T issue Res., 1975, 18, 1.formed amorphous minerals, amorphous calcium phosphate 22 J. D. Currey, T he Mechanical Adaptations of Bones, Princeton and calcium carbonate, are used for structural purposes. Both University Press, Princeton, 1984. these phases need to be stabilized, and they are relatively soft 23 F. Lipmann, Sedimentary Carbonate Minerals, Springer Verlag, compared to their crystalline counterparts.One possibility is Berlin, 1973. that organisms benefit from their isotropic properties. 24 D. Carlstrom, Biol. Bull., 1963, 125, 441. The strive towards isotropy may be a common theme in the 25 H. A. Lowenstam and D. P. Abbot, Science, 1975, 188, 363. 26 A. Hall and J. D. Taylor, MineralMag., 1971, 38, 521. design strategies of many biological materials, and in particular 27 J.R. Clarkson, T. J.Price and C. J. Adams, J. Chem. Soc., Faraday those that are required to fulfil more general functions. The T rans., 1992, 88, 2423. advantages of constructing materials that are more isotropic 28 M. E. Marsh and R. L. Sass, Science, 1980, 208, 1262. have also been recognized by the designers of synthetic com- 29 V. J. Laraia, M. Aindow and A.H. Heuer,Mater. Res. Soc. Symp. posite materials.111 We believe that one potentially promising Proc., 1990, 174, 117. avenue of research in materials science is to reveal some of the 30 M. Sarikaya, J. Liu and I. A. Aksay, in Biomimetics: Design and Processing of Materials, ed. M. Sarikaya and I. A. Aksay, strategies used by organisms to produce more isotropic com- American Institute of Physics, New York, 1995, p. 35. posite materials out of highly anisotropic building blocks. 31 P. Hartman and W. G. Perdok, Acta Crystallogr., 1955, 8, 49. Organisms appear to have had to solve this and many other 32 B. Lawn, Fracture of Brittle Solids, Cambridge University Press, problems relating to their structural materials during the 550 Cambridge, 1993, ch. 6, p. 143. million years of on-the-job testing.Some of the strategies used 33 Y. H. Han, H. Li, T. Y. Wong and R. C. Brad, J. Am. Ceram. Soc., and solutions derived may well haveractical applications in 1991, 74, 3129. 34 M. J.Weedon and P. D. Taylor, Biol. Bull., 1995, 188, 281. the world of synthetic composite materials. 35 E. M. Landau, R. Popovitz-Biro, M. Levanon, L. Leiserowitz, M. Lahav and J.Sagiv,Mol. Cryst. L iq. Cryst., 1986, 134, 323. We thank Joanna Aizenberg, Shira Albeck, Amir Berman, Vivi 36 S. Mann, B. R. Heywood, B. Rajam and J. D. Birchall, Nature Ziv, Giuseppe Falini, Revital Vaknin-Fein and Elia Beniash, (L ondon), 1988, 334, 692. all of whom were our partners in these studies. Fig. 1 and 13 37 B. R. Heywood and S. Mann, Chem.Mater., 1994, 6, 311. 38 L. Addadi, J. Moradian, E. Shay, N. Maroudas and S. Weiner, are from the collection of the late H. A. Lowenstam. These Proc. Natl. Acad. Sci. USA, 1987, 84, 2732. studies were supported by grants from the United States–Israel 39 L. Addadi and S. Weiner, Proc. Natl. Acad. Sci. USA, 1985, 82, Binational Science Foundation and the US Public Health 4110. Service (grant no. DE06954). S.W.holds the I. W. Abel 40 Y. Kitano, N. Kanamori and A. Tokuyama, Am. Zool., 1969, Professorial Chair of Structural Biology, and L. A. the Patrick 9, 681. 41 K. M. Wilbur and A. M. Bernhardt, Biol. Bull., 1984, 166, 251. A. Gorman Professorial Chair of Biological Ultrastructure. 42 R. Giles, S. Manne, S. Mann, D. E. Morse, G. D. Stucky and P. K. Hansma, Biol. Bull., 1995, 188, 8. 43 G. Bevelander and H. Nakahara, Calcif. T issue Res., 1969, 4, 101. References 44 H. Nakahara and M. Kakei, Bull. Yosai Dent. Univ., 1983, 12, 1. 1 S. A. Wainwright, W. D. Biggs, J. D. Currey and J. M. Gosline, 45 S. Weiner and W. Traub, Philos. T rans. R. Soc. L ondon, Ser. B, Mechanical Design in Organisms, Princeton University Press, 1984, 304, 421. Princeton, 1976. 46 S. Weiner and W.Traub, FEBS L ett., 1980, 111, 311. 2 J. D. Birchall, in Biomineralization. Chemical and Biological 47 S.Weiner, Y. Talmon and W. Traub, Int. J. Biol. Macromol., 1983, Perspectives, ed. S. Mann, J. Webb and R. J. P. Williams, VCH, 5, 325. Weinheim, 1989, p. 491. 48 G. Falini, S. Albeck, S. Weiner and L. Addadi, Science, 1996, 3 J. F. V. Vincent, Structural Biomaterials, Macmillan Press, 271, 67.London, 1982. 49 A. M. Belcher, X. H. Wu, R. J. Christensen, P. K. Hansma, 4 A. Hiltner, J. J. Cassidy and E. Baer, Ann. Rev. Mater. Sci., 1985, G. D. Stucky and D. E. Morse, Nature (L ondon), 1996, 381, 56. 15, 455. 50 L. G. Baas-Becking and E. W. Galligher, J. Phys. Chem., 1931, 5 S. A. Tyler and E. S. Barghoorn, Science, 1954, 119, 606. 35, 467. 6 H. A. Lowenstam and S.Weiner, On Biomineralization, Oxford 51 E. A. Minchin, Clathrinidœ. Quart. J.Micr. Sci., 1898, 40, 4, 469. University Press, New York, 1989. 52 K. Okazaki and S. Inoue, Dev. Growth Dier., 1976, 18, 413. 7 M. F. Glaessner, T he Dawn of Animal L ife, Cambridge University 53 J. Aizenberg, J. Hanson, T. H. Koetzle, L. Leiserowitz, S. Weiner Press, Cambridge, 1985. and L. Addadi, Chem. Eur.J., 1995, 1, 414. 8 M. Magaritz, W. T. Holser and J. L. Kirschvink, Nature (L ondon), 54 S. Mann and N. H. C. Sparks, Proc. R. Soc. L ondon, Ser. B, 1988, 234, 441. 1986, 320, 258. J. Mater. Chem., 1997, 7(5), 689–702 70155 A. Berman, D. J. Ahn, A. Lio, M. Almeron, A. Reichert and 85 Q. L. Feng, F. Z. Cui and H. D. Li, Biomimetics, 1996, in press. D. Charych, Science, 1995, 269, 515. 86 S. W. Wise, Eclogae Geol., 1970, 63, 775. 56 W. K. Burton, N. Cabrerra and F. C. Franck, Philos. T rans. 87 S. Weiner and W. Traub, in Structural Aspects of Recognition and R. Soc. L ondon, Ser. A, 1951, 243, 299. Assembly in Biological Macromolecules, ed. M. Balaban, 57 L. Addadi, Z. Berkovitch-Yellin, I. Weissbuch, J. van Mil, J. L. Sussman, W. Traub and A.Yonath, Balaban ISS, Rehovot, L. J. W. Shimon, M. Lahav and L. Leiserowitz, Angew. Chem., Philadelphia, 1981, p. 467. Int. Ed. Engl., 1985, 24, 466. 88 A. P. Jackson, J. F. V. Vincent and R. M. Turner, Proc. R. Soc. 58 S. Mann, Struct. Bonding, 1983, 54, 125. L ondon, Ser. B, 1988, 234, 415. 59 H. Silyn-Roberts and R. M. Sharp, Proc. R. Soc. L ondon, Ser. B, 89 J. D. Currey, Proc.R. Soc. L ondon, Ser. B, 1977, 196, 443. 1986, 227, 303. 90 M. Sarikaya, K. E. Gunnison, M. Yasrebi and I. A. Aksay, Mater. 60 M. Ogilvie, Philos. T rans. R. Soc. L ondon, Ser. B, 1896, 187, 83. Res. Soc. Symp. Proc., 1990, 174, 109. 61 G. L. Decker and W. J. Lennarz, Dev. Biol., 1988, 126, 433. 91 V. J. Laraia and A. H. Heuer, Mater. Res. Soc. Symp. Proc., 1990, 62 K. Okazaki and S. Inoue, Dev. Growth Dier., 1976, 18, 413. 174, 125. 63 R. B. Emlet, Biol. Bull., 1982, 163, 264. 92 S. Weiner, T. Arad, V. Ziv and W. Traub, in Chemistry and 64 G. Donnay and D. L. Pawson, Science, 1969, 166, 1147. Biology of Mineralized T issues, ed. H. Slavkin and P. A. Price, 65 K. Okazaki, Embryologia, 1960, 5, 283. Excerpta Medica, Amsterdam, 1992, p. 93. 66 K. Markel, U. Roser and M. Stauber, Zoomorphology, 1989, 109, 93 R. A. Robinson, J. Bone Joint Surg. A, 1952, 34, 389. 79. 94 S. Weiner and W. Traub, FEBS L ett., 1986, 206, 262. 67 A. Berman, L. Addadi and S. Weiner, Nature (L ondon), 1988, 95 D. E. Birk, E. I. Zycband, D. A. Winkelmann and R. L. Trelstad, 331, 546. Proc. Natl. Acad. Sci. USA, 1989, 86, 4549. 68 A. Berman, L. Addadi, A. Kvick, L. Leiserowitz, M. Nelson and 96 H. C. Anderson, J. Cell Biol., 1969, 41, 59. S.Weiner, Science, 1990, 250, 664. 97 D. M. Kohler, M. A. Crenshaw and A. L. Arsenault,Matrix Biol., 69 J. Aizenberg, S. Albeck, S. Weiner and L. Addadi, J. Cryst. 1994, 14, 543. Growth, 1994, 142, 156. 98 J. Christoersen and W. J. Landis, Anat. Rec., 1991, 230, 435. 70 A. Berman, J. Hanson, L. Leiserowitz, T. Koetzle, S. Weiner and 99 S. F. Jackson, Proc. R. Soc. L ondon, Ser. B, 1957, 146, 270. L. Addadi, Science, 1993, 259, 776. 100 W. Traub, T. Arad and S.Weiner, Matrix, 1992, 12, 251. 71 H. P. Klug and L. E. Alexander, in X-Ray Diraction Procedures, 101 W. Traub, T. Arad and S.Weiner, Conn. T iss. Res., 1992, 28, 99. Wiley Interscience, New York, 1974, 2nd edn., ch. 9. 102 A. L. Arsenault, Calcif. T issue Int., 1988, 43, 202. 72 J. Aizenberg, L. Addadi, J. Hanson, L. Leiserowitz, T. F. Koetzle 103 A. de Ricqle�s, F. J. Meunier, J. Castanet and H. Francillon- and S.Weiner, FASEB J., 1995, 9, 262. Vieillot, in Bone, ed. B. K. Hall, CRC Press, Boca Raton, FL, 73 T. Vaknin-Fein, MSc Thesis, Feinberg Graduate School, 1990, vol. 3, p. 1. Weizmann Institute of Science, 1996. 104 I. R. H. Kramer, Br. Dental J., 1951, 91, 1. 74 P. L. O’Neill, Science, 1981, 213, 646. 105 S. Weiner and W. Traub, FASEB J., 1992, 6, 879. 75 J. Aizenberg, J. Hanson, T. F. Koetzle, S. Weiner and L. Addadi, 106 V. Ziv, I. Sabanay, T. Arad, W. Traub and S.Weiner,Micros. Res. J. Am. Chem. Soc., in press. T echnol., 1996, 33, 203. 76 Silicon and Siliceous Structures in Biological Systems, ed. 107 V. Ziv, H. D. Wagner and S.Weiner, Bone, 1996, 18, 417. T. L. Simpson and B. E. Volcani, Springer Verlag, New York, 108 J. G. Carter and G. R. Clark, in Mollusks. Notes for a Short 1981. Course, ed. T. W. Broadhead, University of Tennessee, Knoxville, 77 T. Pobeguin, Ann. Sci. Nat. Bot., Ser. II, 1954, 15, 29. TN, 1985, p. 50. 78 H. A. Lowenstam, Bull. Mar. Sci., 1989, 45, 243. 109 H. A. Lowenstam and S. Weiner, Proc. Natl. Acad. Sci. USA, 79 E. Beniash, J. Aizenberg, L. Addadi and S. Weiner, Philos. T rans. 1992, 89, 10573. R. Soc. L ondon, in press. 110 C. C. Harrison, Phytochem., 1996, 41, 37. 80 K. M. Wilbur and N.Watabe, Ann. N.Y. Acad. Sci., 1963, 109, 82. 111 J. G. Parkhouse and A. Kelly, Proc. R. Soc. L ondon, Ser. A, 1995, 81 W. U. Berthold, Naturwisschenschaften, 1976, 63, 196. 451, 737. 82 T. R. Waller, Smithsonian Contrib. Biol., 1980, 313, 1. 112 W. Traub, T. Arad and S. Weiner, Proc. Natl. Acad. Sci. USA, 83 S.Weiner, CRC Crit. Rev. Biochem., 1986, 20, 365. 1989, 86, 9822. 84 S. Manne, C. M. Zaremba, R. Giles, L. Huggins, D. A. Walters, A. Belcher, D. E. Morse, G. D. Stucky, J. M. Didymus, S. Mann and P. K. Hansma, Proc. R. Soc. L ondon, Ser. B, 1994, 256, 17. Paper 6/04512J; Received 28th June, 1996 702 J. Mater. Chem., 1997, 7(5), 689–7

ISSN:0959-9428    

DOI:10.1039/a604512j

出版商:RSC    

年代:1997

数据来源: RSC

摘要:

MATERIALS CHEMISTRY COMMUNICATION Magneto-optical properties of nanostructured iron K. E. Gonsalves,a,b G. Carlson,*a M. Benaissa,c M. Jose-Yacama�n,c D. Y. Kimd and J. Kumard aPolymer Program, Institute of Materials Science, University of Connecticut, Storrs, CT 06269, USA bDepartment of Chemistry, University of Connecticut, Storrs, CT 06269, USA cInstituto de Fisica, Universidad Nacional Autonoma de Mexico, Apdo.Postal 20-364, C.P. 01000,Mexico, D.F.,Mexico dCenter for Advanced Materials, Departments of Physics and Chemistry, University ofMassachusetts, L owell, L owell, MA 01864, USA Particles of a-iron of 2–10 nm size have been prepared by chemical synthesis, stabilized and dispersed into a polymer matrix. The composite was characterized using high-resolution transmission electron microscopy.The optical Faraday rotation of the polymer composite, measured at 633 nm, was approximately 1.2° T-1 . The Verdet constant measured by the Faraday rotation of the polymer composite at 633 nm is 0.36 min Oe-1 cm-1. Nanostructured materials1 are those with a grain or domain size of less than 100 nm. This significant reduction in size can produce dramatic changes in the behaviour of the material.Because of its high coercivity,2 iron is often used in magnetic applications. By using nanostructured iron, the eciency of magneto-caloric refrigeration and magnetic data storage can be improved.3 Conventional iron, with larger (mm-sized) domains, is usually prepared by physical routes such as arc melting.4 Nanostructured materials are also prepared by physical methods, including sputtering and molecular beam epitaxy (MBE).5 However, there has recently been much interest in the chemical synthesis of nanostructured materials.This approach has the advantages of improved mixing, better control of stoichiometry, and tailored synthesis by assembly of atomic or molecular precursors.6 Iron nanoparticles, either as a solid or incorporated into an amorphous matrix, have been shown to exhibit superparamagnetism at low temperatures.7 Owing to the small size of the nanoparticles the cooperative behaviour of the spins does not lead to ferromagnetism.At room temperature the composite material has paramagnetic behaviour. It is of interest to examine the possibility of using nanoparticles of iron and iron oxide in a polymer matrix for use as a Faraday rotator.Ironcontaining ferrites such as yttrium iron garnet and bismuth iron garnet are already used as Faraday rotators for optical isolator applications.8 The possibility of a polymeric Faraday rotator using nanoparticles which can be moulded or cast into rods without the expense of crystal growth and polishing is very attractive.We have therefore made measurements of the Verdet constant of nitrogen-passivated iron particles in a poly(methyl methacrylate) (PMMA) matrix at 633 nm and 1.3 mm. Nanostructured iron powder was prepared by thermal decomposition9 of Fe(CO)5 in decalin solution. During the decomposition and coolingof the mixture, anhydrous ammonia gas was passed through the solution.The magnetic moment of the iron powder was 60 emu g-1 (cgs, Gaussian) at room temperature and 120 emu g-1 at 10 K. To prepare the polymer–matrix composite, approximately 500 mg of the powder and 2 ml of dodecanethiol were sealed in a flask under nitrogen and sonicated in a cleaning bath for 2 h.The thiol was then decanted and methyl methacrylate was added, along with azoisobutyronitrile, an initiator. This mixture was polymerized by heating at 72°C for 2 h. The composite sample was examined using a JEOL-4000EX electron microscope with an accelerating voltage of 400 kV and a point-to-point resolution of approximately 1.7 A°. HRTEM images were obtained at optimum (Scherzer) defocus.Computer filtering was used to improve the images. The HRTEM characterization (Fig. 1) of the particles in the PMMA matrix showed that the size distribution is narrow. Symmetry analysis and measurement of the lattice fringe spacings (Fig. 2) in the HRTEM image indicate that the particle is a-Fe (bcc) observed along the 001 zone axis.Particles show flat surfaces, terminating with {110} facets. This seems reasonable, since the (110) planes are most dense in a bcc lattice, which makes them very stable. Surprisingly, the particles were not oxidized. XPS results indicate that nitrogen is present, with a 1552 iron5nitrogen ratio. It seems likely that nitrogen is on the particle surface, preventing oxidation.However, this is dicult to confirm through TEM because of nitrogen’s weak scattering intensity and further studies using NMR and EXAFS are under way. The magnetic susceptibility of the composite material was Fig. 1 High-resolution transmission electron microscope image of nanostructured iron/PMMA composite J. Mater. Chem., 1997, 7(5), 703–704 703Fig. 2 Computer filtered image of the iron nanoparticle. Black corresponds to atomic columns. Enlargement of nano-Fe HRTEM image. measured in aGuoy susceptibility balance at room temperature and found to be 1.2×10-5 esu. The PMMA matrix has a diamagnetic contribution to the Verdet constant (Vd) while the iron nanoparticles at room temperature have a paramagnetic contribution (Vp). The Verdet constant at fixed temperature T and wavelength l is given by V (l,T )=Vp(l,T )+Vd(l,T ) Magneto-optical experiments were conducted on the samples in the transmission geometry.A polarized laser beam at the desired wavelength (0.633 or 1.3 mm) was transmitted through a sample of the composite material and was collinear with the direction of the applied magnetic field. The magnetic field strength was varied from 0 to 1.5 T.The output polarizer P was oriented at an angle of 45° to the direction of the input polarization. Therefore, if the polarizer is aligned in the direction of the input polarization (Iout=Io cos2 h), the transmission without an applied magnetic field is half the output intensity value. The application of a magnetic field results in a change of the output intensity which is proportional to the rotation of the plane of polarization of the output beam with respect to the input polarization.For small rotations, the change in relative output intensity is equal to the rotation of the plane of polarization in radians. The slope of this curve gives the value of the Verdet constant of the composite material.At 633 nm the transmission of a 200 mm sample is 5%. The losses may be due to a combination of absorption and scattering and no attempt has been made to minimize them at this stage. Fig. 3 shows the magnitude of rotation of the plane of polarization of the output beam with the applied field at 633 nm. The best linear fit to the data passing through the origin gives a slope of 1.2° T-1.The scatter in the data points give an uncertainty of 10% in the slope and consequently the Verdet constant. The Verdet constant calculated from the slope in units most often quoted in the literature is 0.36 min Oe-1 cm-1. The transmission at 1.3 mm is about 40% in a 500 mm thick sample. The value measured at 1.3 mm from a 500 mm thick sample is 0.1 min Oe-1 cm-1.The drop in Verdet constant at 1.3 mm is theoretically expected. In paramagnetic materials at wavelength far from resonance the Verdet constant decreases as 1/l2. 704 J. Mater. Chem., 1997, 7(5), 703–704 Fig. 3 Faraday rotation as a function of magnetic field It is interesting to note that at room temperature the Verdet constant at 1.3 mm is considerably larger than the commercially available Faraday rotator glass10 Hoya FR-5.A uniform dispersion was obtained and phase separation can be avoided by surface functionalization. This approach can be applied to other ferromagnetic or ferrite nanoparticles. Also, nanoparticulate rare-earth-metal trifluorides11 may be investigated. The polymeric–Fe composite material has not been optimized for its Faraday rotation and scattering losses.These issues need to be further investigated before the polymeric composite material can emerge as an inexpensive alternative to existing inorganic single crystals or glasses. Nevertheless the values of Verdet constant possible in polymeric compositeare encouraging.K. E. Gonsalves wishes to acknowledge partial support of this work by ONR grant N0014-94-1-0833 and NSF grant INT- 9503854. References 1 H. Gleiter, Adv. Mater., 1992, 4, 474. 2 R. Birringer and H. Gleiter, in Encyclopedia of Materials Science and Engineering, ed. R. W. Cahn, Pergamon Press, Oxford, 1988, suppl. vol. 1, p. 339. 3 Y. Teng and B. Li, J. Funct. Mater., 1994, 25, 116; R.Sessoli, D. Gatteschi, A. Caneschi and M. A. Novak, Nature (L ondon), 1993, 365, 141. 4 H. H. Stadelmaier and E.-T. Henig, J. Mater. Eng. Performance, 1992, 1, 167. 5 M. N. Baibich, J. M. Broto, A. Fert, F. Nguyen van Dau,F. Petro, P. Etienne, G. Creuzet, A. Friederich and J. Charles, Phys. Rev. L ett., 1988, 61, 2472. 6 K. E. Gonsalves, ed., 1996, Nanotechnology: Molecularly Designed Materials, ed. K. E. Gonsalves, ACS Symp. Ser. 622, 1996. 7 C. P. Bean and J. D. Livingston, J. Appl. Phys., 1959, 30, 120. 8 M. Okada, S. Katayama and K. Tominaga, J. Appl. Phys., 1991, 69, 3566; K. Matsumoto, S. Sasaki and K. Haraga, J. Appl. Phys., 1992, 71, 2467. 9 K. E. Gonsalves, US Pat., 4842 641, 1989. 10 J. A. Davis and R. M. Bunch, Appl. Opt., 1984, 23, 633; Handbook of L aser Science and T echnology, ed. M. J.Weber, CRC Press, Boca Raton, FL, 1988, vol. IV. 11 C. Leycuras, H. LeGall, M. Guillot and A. Marchand, J. Appl. Phys., 1984, 55, 2161. Communication 7/00899F; Received 7th February, 1997

ISSN:0959-9428    

DOI:10.1039/a700899f

出版商:RSC    

年代:1997

数据来源: RSC

摘要:

Synthesis and structural study of a novel nonlinear optical material: the tolane derivative ethyl 2-(4-benzyloxyphenylethynyl )-5-nitrobenzene-1-carbamate Midori Kato,*a Kimiko Kobayashi,b Masaaki Okunaka,a Nami Sugita,a Masashi Kiguchia and Yoshio Taniguchia aAdvanced Research L aboratory, Hitachi, L td., Hatoyama, Saitama 350-03, Japan bT he Institute of Physical and Chemical Research (RIKEN), 2-1 Hirosawa, Wako-shi, Saitama 351-01, Japan We have synthesized the tolane derivative ethyl 2-(4-benzyloxyphenylethynyl)-5-nitrobenzene-1-carbamate 1, a promising material for nonlinear optics.Compound 1 has several crystal forms and we have characterized their structural properties by X-ray diraction. One of the three forms, which is crystallized from a strong polar solvent, has a non-centrosymmetric structure and is ca.ten times as strong a second-harmonics generator as meta-nitroaniline. A hydrogen bond is found to perform an important role in forming the non-centrosymmetric geometry, which is necessary for a nonlinear optical crystal. A great deal of work has been performed to make new ecient dehyde and benzyl bromide with malonic acid.10 The acid was treated with bromine, followed by elimination of HBr in two materials for nonlinear optics.Organic materials are expected to have high nonlinear optical eciency.1,2 The large second- steps, which resulted in the formation of 4-benzyloxyphenylacetylene. 11 On the other hand, ethyl 2-bromo-5-nitrobenzene-1- order molecular nonlinearity of most organic materials originates from intramolecular charge transfer.3–5 This interaction carbamate was prepared by the reaction of 4-bromo-3-aminonitrobenzene with ethyl chloroformate.12 Ethyl 2-bromo-5- is enhanced by donor and acceptor substituents bridged by molecules with a p-electron system.nitrobenzene-1-carbamate and 4-benzyloxyphenylacetylene were reacted using CuI and PdCl2(PPh3)2 in NEt3 to form To achieve ecient nonlinear optical devices, both high hyperpolarizability and an appropriate crystal structure is ethyl 2-(4-benzyloxyphenylethynyl)-5-nitrobenzene-1-carbamate 1. The obtained material has a yellow colour.The needed. Although a non-centrosymmetric structure is required for this purpose, molecules with a donor–acceptor system tend absorption spectrum of 1 recrystallized from 2-ethoxyethanol is shown in Fig. 1. The cut-o wavelength is 523 nm, which to crystallize in a centrosymmetric structure. Tolane† derivatives are known to have high hyperpolariz- was shortened by the purification. ability due to the large charge transfer between the two benzene rings bridged by the central triple bond, in spite of a low cut- Dierential scanning calorimetry o wavelength.6 As is often the case with organic donor– The thermal analysis was performed with a Perkin-Elmer DSC acceptor systems, tolane derivatives also tend to have a centro- 7.The cooling and heating rates were 5°C min-1. After DSC symmetric structure and it turns out that the second-order measurement the IR spectrum or the DSC curve of the sample nonlinearity vanishes.7 More than 150 tolane derivatives have was measured to confirm that the sample had not decomposed.been synthesized and their SHG eciencies in powder form have been tested by a second-harmonic wave with the evan- Crystal structure determination escent wave (SHEW) technique,8 some of which have already been reported.9 Among them we picked out ethyl 2-(4-benzyl- The single crystals for the X-ray diraction study were grown oxyphenylethynyl)-5-nitrobenzene-1-carbamate 1 (4-nitro-2- in solution by the cooling method.Compound 1 shows polyethoxyamido- 4¾-benzyloxytolane) because it was found to be a morphism. As far as we know, compound 1 has three types of promising material during the preliminary investigations. It crystal structure (designated a, b and c for simplicity).has a benzyloxy group as a donor and a nitro group as an To crystallize the a-form crystal, purified 1 was dissolved in acceptor. Using dierential scanning calorimetry (DSC), we 2-ethoxyethanol at 40°C until the solution was almost satufound that 1 has more than two forms of crystal structure. We identified the 2 types of structure by single-crystal X-ray diraction and discuss here the condition of the appropriate crystal for nonlinear optics.CH2O CC NHCO2Et NO2 1 Experimental Synthesis The synthesis of compound 1 is briefly described. 4- Benzyloxycinnamic acid was prepared by reaction of 4-benzyloxybenzaldehyde, which was prepared from 4-hydroxybenzal- Fig. 1 The absorption spectrum of the a-form crystal † IUPAC name: diphenylacetylene. J.Mater. Chem., 1997, 7(5), 705–711 705Table 1 Crystallographic data a- and b-form crystals sample name a-form b-form molecular formula C24H20N2O5 C24H20N2O5 molecular weight 416.42 416.42 crystal system monoclinic monoclinic space group Cc (No. 9) P21/c (No. 14) unit cell dimensions a/A° 25.827(5) 10.514(1) b/A° 4.947(2) 42.385(4) c/A° 21.089(8) 14.185(2) b/degrees 127.78(4) 90.47(1) unit cell volume U (A° 3) 2129.5(6) 6321(1) Z 4 12 D(mg m-3) 1.299 1.313 no.of observed reflections 2614 7829 no. of refined parameters 361 1039 R 0.056 0.087 computer and programs FACOM M-1800, UNICS-III program system Table 2 Selected bond lengths of the a-form crystala bond length/A° bond length/A° O(17)–C(16) 1.335(0.006) O(17)–C(18) 1.466(0.009) O(20)–C(16) 1.188(0.008) O(22)–N(21) 1.230(0.007) O(23)–N(21) 1.202(0.011) O(24)–C(12) 1.371(0.009) O(24)–C(25) 1.398(0.012) N(15)–C(3) 1.402(0.006) N(15)–C(16) 1.359(0.008) N(15)–H(N15) 0.946(0.057) N(21)–C(5) 1.440(0.010) C(1)–C(2) 1.423(0.010) C(1)–C(8) 1.184(0.010) C(2)–C(3) 1.426(0.011) C(2)–C(7) 1.409(0.008) C(3)–C(4) 1.380(0.010) C(4)–C(5) 1.395(0.007) C(5)–C(6) 1.417(0.012) C(6)–C(7) 1.359(0.011) C(8)–C(9) 1.443(0.010) C(9)–C(10) 1.406(0.008) C(9)–C(14) 1.376(0.013) C(10)–C(11) 1.376(0.012) C(11)–C(12) 1.404(0.013) C(12)–C(13) 1.365(0.008) C(13)–C(14) 1.397(0.011) C(18)–C(19) 1.468(0.011) C(25)–C(26) 1.528(0.010) C(26)–C(27) 1.375(0.009) C(26)–C(31) 1.344(0.014) C(27)–C(28) 1.380(0.012) C(28)–C(29) 1.359(0.016) C(29)–C(30) 1.352(0.011) C(30)–C(31) 1.380(0.011) aE.s.d.s shown in parentheses. crystal structures were determined by direct methods using MULTAN78, and were refined by use of the block-diagonal least-square method.Fig. 2 (a) DSC curve of compound 1. Starting from the a-form crystal, Atomic coordinates, thermal parameters, and bond lengths the sample was heated to 150°C and kept at that temperature for and angles have been deposited at the Cambridge Crys- 5 min.Then it was cooled to room temperature. The arrows show the tallographic Data Centre (CCDC). See Information for direction of heating and cooling. (b) DSC curve starting with the b- Authors, J. Mater. Chem., 1997, Issue 1. Any request to the form crystal. The sample was kept at 150°C for 5 min. (c) DSC curve CCDC for this material should quote the full literature citation starting with the b-form crystal.The sample was kept at 135°C for 5 min, then cooled down. and the reference number 1145/29. Linear and nonlinear optical measurements rated and was then cooled slowly at a rate of 0.1–0.2 °C h-1. After the solution had reached room temperature, crystals The UV–visible absorption spectra were measured with a HITACHI spectrophotometer U-3400.were picked from the solution and rinsed with a solvent to clean the crystal surfaces. The acquired crystal had a The directions of the optical axes of the crystal were defined by the conoscopic figure with the polarizing microscope at a long, narrow shape, with dimensions of approximately 1×3×0.4 mm. The crystal had partially plane and parallel wavelength of 532 nm.The refractive indices were measured by the observation of Becke’s line. To fit the Sellmeier equation, surfaces. This area was about 1×1 mm in size. To obtain the b-form crystal, the same procedure was measurements at three wavelengths (532, 589 and 633 nm) were performed. All the optical measurements were performed performed but ethyl acetate was used as a solvent. The b-form crystals easily formed as a large parallelogram platelet crystal only on the a-form crystals.Nonlinear optical properties were evaluated using the of dimensions 10×10×1 mm. The c-form crystal appears to be unstable, because it only SHEW technique. Details of this technique are described in ref. 8. The compound was crushed into a fine powder and appears during the thermal processes as mentioned in the next section.It was not obtained at room temperature. pushed into contact with a surface of a total reflection prism made of rutile. The fundamental wave of aQ-switched Nd:YAG Cell parameters and intensity data for the a-form crystals were derived from measurements on a Rigaku AFC-4 four- laser (l=1064 nm) was guided into the prism and was totally reflected at the boundary.The reflected second-harmonics circle diractometer and the data for the b-form crystals were from an Enraf-Nonius CAD-4 diractometer. Molecular and (SH) were detected and their power was compared with that 706 J. Mater. Chem., 1997, 7(5), 705–711Table 4 Selected bond lengths of the b-form crystala of meta-nitroaniline (m-NA) measured under the same conditions. bond length/A° bond length/A° O(17)A–C(16)A 1.336(0.007) O(17)A–C(18)A 1.502(0.008) Results and Discussion O(20)A–C(16)A 1.196(0.007) O(22)A–N(21)A 1.216(0.007) O(23)A–N(21)A 1.218(0.007) O(24)A–C(12)A 1.380(0.007) Dierential scanning calorimetry (DSC) analysis shows how O(24)A–C(25)A 1.410(0.008) N(15)A–C(3)A 1.395(0.007) the a-, b- and c-forms appear.Fig. 2(a) shows the DSC curve N(15)A–C(16)A 1.381(0.007) N(15)A–H(N15)A 0.968(0.054) produced when the a-form crystal was heated.The crystal was N(21)A–C(5)A 1.502(0.007) C(1)A–C(2)A 1.432(0.008) heated to 150 °C, the temperature maintained for 5 min and C(1)A–C(8)A 1.176(0.008) C(2)A–C(3)A 1.418(0.008) then cooled to 30°C. The a-form crystal melts at 143.1 °C. In C(2)A–C(7)A 1.383(0.008) C(3)A–C(4)A 1.388(0.008) the cooling process a small exothermic peak was observed at C(4)A–C(5)A 1.394(0.008) C(5)A–C(6)A 1.347(0.008) C(6)A–C(7)A 1.383(0.008) C(8)A–C(9)A 1.464(0.008) 63.3 °C after the solidification.Since the final crystal obtained C(9)A–C(10)A 1.401(0.009) C(9)A–C(14)A 1.374(0.008) after the entire process proved to be the b-form crystal, this C(10)A–C(11)A 1.376(0.009) C(11)A–C(12)A 1.403(0.009) small exothermic peak can be regarded as a phase transition C(12)A–C(13)A 1.381(0.009) C(13)A–C(14)A 1.402(0.008) to the b-form crystal.It might be thought that the b-form C(18)A–C(19)A 1.407(0.010) C(25)A–C(26)A 1.517(0.008) crystal is transformed from the a-form crystal at 63.3 °C, but C(26)A–C(27)A 1.375(0.009) C(26)A–C(31)A 1.377(0.009) the DSC curve in Fig. 2(b) shows another explanation. C(27)A–C(28)A 1.375(0.009) C(28)A–C(29)A 1.367(0.011) C(29)A–C(30)A 1.365(0.010) C(30)A–C(31)A 1.392(0.009) Fig. 2(b) illustrates the heating and cooling processes starting O(17)B–C(16)B 1.351(0.007) O(17)B–C(18)B 1.522(0.008) from the b-form crystal. When the b-form crystal was heated, O(20)B–C(16)B 1.194(0.007) O(22)B–N(21)B 1.216(0.007) a transition occurred at 65.5 °C which was indicated by a small O(23)B–N(21)B 1.230(0.006) O(24)B–C(12)B 1.366(0.007) endothermic peak.Another transition was observed at 135.5 °C O(24)B–C(25)B 1.440(0.007) N(15)B–C(3)B 1.396(0.007) and then the crystal melted at the melting point of the a-form N(15)B–C(16)B 1.368(0.007) N(15)B–H(N15)B 0.934(0.055) crystal; the second transition at 135.5 °C is therefore the N(21)B–C(5)B 1.500(0.007) C(1)B–C(2)B 1.431(0.008) C(1)B–C(8)B 1.186(0.008) C(2)B–C(3)B 1.416(0.008) transition to the a-form crystal, and thus the first transition C(2)B–C(7)B 1.403(0.008) C(3)B–C(4)B 1.381(0.008) must be considered as the transition to another form, which is C(4)B–C(5)B 1.386(0.008) C(5)B–C(6)B 1.370(0.008) the c-form.The IR spectra at 80°C also showed that the C(6)B–C(7)B 1.381(0.008) C(8)B–C(9)B 1.462(0.008) crystal in the intermediate is very dierent from the a-form C(9)B–C(10)B 1.392(0.008) C(9)B–C(14)B 1.391(0.008) crystal, that is, the c-form crystal.Although the detailed C(10)B–C(11)B 1.364(0.008) C(11)B–C(12)B 1.380(0.008) structure of the c-form has not been investigated as yet, it is C(12)B–C(13)B 1.401(0.008) C(13)B–C(14)B 1.398(0.008) C(18)B–C(19)B 1.284(0.012) C(25)B–C(26)B 1.513(0.008) expected that it is little dierent from the b-form crystal C(26)B–C(27)B 1.357(0.009) C(26)B–C(31)B 1.400(0.009) because the IR spectrum of the c-form crystal was similar to C(27)B–C(28)B 1.386(0.009) C(28)B–C(29)B 1.348(0.009) that of the b-form crystal.Thus the dierence between the b- C(29)B–C(30)B 1.375(0.010) C(30)B–C(31)B 1.396(0.009) and c-forms is probably a small conformational change in the O(17)C–C(16)C 1.357(0.007) O(17)C–C(18)C 1.526(0.008) benzyloxy or carbamate substituents, or a dierent rotational O(20)C–C(16)C 1.213(0.007) O(22)C–N(21)C 1.226(0.007) angle around the COC axis.O(23)C–N(21)C 1.224(0.007) O(24)C–C(12)C 1.368(0.007) O(24)C–C(25)C 1.435(0.008) N(15)C–C(3)C 1.387(0.007) On the other hand, the enthalpy of the smaller exothermic N(15)C–C(16)C 1.341(0.007) N(15)C–H(N15)C 1.040(0.056) peak at 63.3 °C in the cooling process is equivalent to that of N(21)C–C(5)C 1.497(0.007) C(1)C–C(2)C 1.419(0.009) the endothermic peak at 65.4 °C in the heating process.C(1)C–C(8)C 1.194(0.009) C(2)C–C(3)C 1.419(0.008) Therefore the crystal formed during cooling from 120.3 to C(2)C–C(7)C 1.405(0.009) C(3)C–C(4)C 1.394(0.008) 65.4 °C was also the c-form crystal.The cooling curve in C(4)C–C(5)C 1.388(0.008) C(5)C–C(6)C 1.375(0.008) Fig. 2(a) is same as that in Fig. 2(b), so that in the cooling C(6)C–C(7)C 1.382(0.009) C(8)C–C(9)C 1.450(0.009) C(9)C–C(10)C 1.405(0.009) C(9)C–C(14)C 1.397(0.008) process the c-form always crystallizes from the melted com- C(10)C–C(11)C 1.355(0.009) C(11)C–C(12)C 1.400(0.009) pound 1 and then transforms to the b-form crystal at 65.4 °C.C(12)C–C(13)C 1.389(0.009) C(13)C–C(14)C 1.404(0.009) Fig. 2(c) shows the possibility of obtaining the a-form crystals C(18)C–C(19)C 1.470(0.010) C(25)C–C(26)C 1.516(0.009) C(26)C–C(27)C 1.367(0.009) C(26)C–C(31)C 1.398(0.009) Table 3 Selected bond angles of the a-form crystala C(27)C–C(28)C 1.377(0.010) C(28)C–C(29)C 1.364(0.010) C(29)C–C(20)C 1.378(0.010)709C(30)C–C(31)C 1.396(0.009) bonds angle (°) bonds angle (°) aE.s.d.s shown in parentheses.C(16)–O(17)–C(18) 115.1(0.5) C(12)–O(24)–C(25) 116.7(0.5) C(3)–N(15)–C(16) 123.9(0.5) C(3)–N(15)–H(N15) 124.2(3.1) from the b-form crystals. When the b-form crystal was heated C(16)–N(15)–H(N15) 111.9(3.1) O(22)–N(21)–O(23) 123.0(0.7) O(22)–N(21)–C(5) 117.5(0.7) O(23)–N(21)–C(5) 119.6(0.5) until the transition to the a-form at 135.5 °C, and then cooled C(2)–C(1)–C(8) 174.4(0.6) C(1)–C(2)–C(3) 119.7(0.5) down before it melted, the a-form remained.The a-form C(1)–C(2)–C(7) 121.5(0.7) C(3)–C(2)–C(7) 118.8(0.7) structure was confirmed by IR spectroscopy and an X-ray N(15)–C(3)–C(2) 117.0(0.6) N(15)–C(3)–C(4) 122.2(0.7) diraction pattern.While the newly-formed a-form crystals C(2)–C(3)–C(4) 120.8(0.5) C(3)–C(4)–C(5) 118.3(0.7) outwardly maintained the large size of the b-form crystals, the N(21)–C(5)–C(4) 118.4(0.7) N(21)–C(5)–C(6) 119.4(0.5) crystals became opaque at the transition, indicating that they C(4)–C(5)–C(6) 122.2(0.6) C(5)–C(6)–C(7) 118.7(0.5) C(2)–C(7)–C(6) 121.3(2.9) C(1)–C(8)–C(9) 177.0(0.8) were no longer single crystals.This is because the structures C(8)–C(9)–C(10) 118.6(0.8) C(8)–C(9)–C(14) 121.2(0.5) of the a- and b-form crystals are very dierent, as mentioned C(10)–C(9)–C(14) 120.2(0.7) C(9)–C(10)–C(11) 119.8(0.8) in the following. C(10)–C(11)–C(12) 119.7(0.6) O(24)–C(12)–C(11) 114.8(0.5) All the obtained parameters of the a- and b-forms of 1 from O(24)–C(12)–C(13) 125.0(0.8) C(11)–C(12)–C(13) 120.2(0.7) the X-ray diraction studies are summarized in Table 1.The C(12)–C(13)–C(14) 120.5(0.8) C(9)–C(14)–C(13) 119.5(0.6) selected bond lengths and the angles of the a-form crystals are O(17)–C(16)–O(20) 124.6(0.5) O(17)–C(16)–N(15) 108.9(0.5) O(20)–C(16)–N(15) 126.4(0.5) O(17)–C(18)–C(19) 107.1(0.7) summarized in Tables 2 and 3, and those of the b-form crystals O(24)–C(25)–C(26) 109.7(0.5) C(25)–C(26)–C(27) 124.1(0.8) are in Tables 4 and 5, respectively.The space group of the a- C(25)–C(26)–C(31) 117.8(0.6) C(27)–C(26)–C(31) 118.2(0.7) form was determined to be monoclinic Cc, which is non- C(26)–C(27)–C(28) 120.2(1.0) C(27)–C(28)–C(29) 120.3(0.7) centrosymmetric, while on the other hand, the b-form crystals C(28)–C(29)–C(30) 119.7(0.8) C(29)–C(30)–C(31) 119.5(1.0) are centrosymmetric, and therefore not SH active. C(26)–C(31)–C(30) 122.1(0.7) For the future study of this potentially SH active material, we characterized the linear optical properties of the a-form aE.s.d.s shown in parentheses.J. Mater.Chem., 1997, 7(5), 705–711 707Table 5 Selected bond angles of the b-form crystala bonds angle (°) bond angle (°) C(16)A–O(17)A–C(18)A 111.4(0.4) C(12)A–O(24)A–C(25)A 118.7(0.5) C(3)A–N(15)A–C(16)A 125.6(0.5) C(3)A–N(15)A–H(N15)A 120.1(3.23) C(16)A–N(15)A–H(N15)A 114.1(3.3) O(22)A–N(21)A–O(23)A 124.5(0.5) O(22)A–N(21)A–C(5)A 117.2(0.5) O(23)A–N(21)A–C(5)A 118.4(0.5) C(2)A–C(1)A–C(8)A 176.7(0.6) C(1)A–C(2)A–C(3)A 118.5(0.5) C(1)A–C(2)A–C(7)A 121.0(0.5) C(3)A–C(2)A–C(7)A 120.5(0.5) N(15)A–C(3)A–C(2)A 118.1(0.5) N(15)A–C(3)A–C(4)A 122.7(0.5) C(2)A–C(3)A–C(4)A 119.3(0.5) C(3)A–C(4)A–C(5)A 117.0(0.5) N(21)A–C(5)A–C(4)A 115.3(0.5) N(21)A–C(5)A–C(6)A 119.8(0.5) C(4)A–C(5)A–C(6)A 124.9(0.5) C(5)A–C(6)A–C(7)A 118.2(0.5) C(2)A–C(7)A–C(6)A 120.2(0.5) C(1)A–C(8)A–C(9)A 178.9(0.6) C(8)A–C(9)A–C(10)A 118.3(0.5) C(8)A–C(9)A–C(14)A 121.1(0.5) C(10)A–C(9)A–C(14)A 120.5(0.6) C(9)A–C(10)A–C(11)A 120.0(0.6) C(10)A–C(11)A–C(12)A 119.0(0.6) O(24)A–C(12)A–C(11)A 113.8(0.5) O(24)A–C(12)A–C(13)A 124.8(0.5) C(11)A–C(12)A–C(13)A 121.4(0.6) C(12)A–C(13)A–C(14)A 118.7(0.5) C(9)A–C(14)A–C(13)A 120.2(0.5) O(17)A–C(16)A–C(20)A 126.0(0.5) O(17)A–C(16)A–N(15)A 108.2(0.5) O(20)A–C(16)A–N(15)A 125.9(0.5) O(17)A–C(18)A–C(19)A 107.1(0.5) O(24)A–C(25)A–C(26)A 107.5(0.5) C(25)A–C(26)A–C(27)A 122.4(0.6) C(25)A–C(26)A–C(31)A 118.3(0.5) C(27)A–C(26)A–C(31)A 119.3(0.6) C(26)A–C(27)A–C(28)A 120.3(0.7) C(27)A–C(28)A–C(29)A 120.6(0.7) C(28)A–C(29)A–C(30)A 119.6(0.6) C(29)A–C(30)A–C(31)A 120.4(0.6) C(26)A–C(31)A–C(30)B 119.8(0.6) C(12)B–N(24)B–C(25)B 117.0(0.4) C(16)B–O(17)B–C(18)B 112.6(0.5) C(3)B–N(15)B–H(N15)B 119.0(0.4) C(3)B–N(15)B–C(16)B 125.7(0.5) O(22)B–N(21)B–O(23)B 124.8(0.5) C(16)B–N(15)B–H(N15)B 115.3(3.4) O(23)B–C(21)B–C(5)B 118.1(0.5) O(22)B–N(21)B–C(5)B 117.1(0.5) C(1)B–C(2)B–C(3)B 120.2(0.5) C(2)B–C(1)B–C(8)B 179.6(0.6) C(2)B–C(2)B–C(7)B 119.9(0.5) C(1)B–C(2)B–C(7)B 119.9(0.5) N(15)B–C(3)B–C(4)B 123.3(0.5) N(15)B–C(3)B–C(2)B 117.1(0.5) C(3)B–C(4)B–C(5)B 117.7(0.5) C(2)B–C(3)B–C(4)B 119.6(0.5) N(21)B–C(5)B–C(6)B 118.4(0.5) N(21)B–C(5)B–C(4)B 116.7(0.5) C(5)B–C(6)B–C(7)B 117.3(0.5) C(4)B–C(5)B–C(6)B 124.9(0.5) C(1)B–C(8)B–C(9)B 176.8(0.6) C(2)B–C(7)B–C(6)B 120.6(0.5) C(8)B–C(9)B–C(14)B 118.9(0.5) C(8)B–C(9)B–C(10)B 121.9(0.5) C(9)B–C(10)B–C(11)B 120.9(0.5) C(10)B–C(9)B–C(14)B 119.2(0.5) C(10)B–C(11)–C(12)B 120.1(0.5) O(24)B–C(12)B–C(13)B 123.7(0.5) O(24)B–C(12)B–C(11)B 115.6(0.5) C(12)B–C(13)B–C(14)B 118.5(0.5) C(11)B–C(12)B–C(13)B 120.7(0.5) O(17)B–C(16)B–O(20)B 124.4(0.5) C(9)B–C(14)B–C(13)B 120.5(0.5) O(20)B–C(16)B–N(15)B 126.4(0.5) O(17)B–C(16)B–N(15)B 109.3(0.5) O(24)B–C(25)B–C(26)B 107.6(0.5) O(17)B–C(18)B–C(19)B 111.1(0.6) C(25)B–C(26)B–C(31)B 119.1(0.7) C(25)B–C(26)B–C(27)B 120.7(0.6) C(26)B–C(27)B–C(28)B 120.3(0.6) C(27)B–C(26)B–C(31)B 120.2(0.5) C(29)B–C(30)B–C(31)B 120.6(0.6) C(27)B–C(28)B–C(29)B 120.6(0.6) C(16)C–O(17)C–C(18)C 115.5(0.4) C(26)B–C(31)B–C(30)B 118.2(0.6) C(3)C–N(15)C–C(16)C 125.9(0.5) C(12)C–O(24)C–C(25)B 117.3(0.5) C(16)C–N(15)C–H(N15)C 121.7(3.5) C(3)C–N(15)C–H(N15)C 111.9(3.1) O(22)C–N(21)C–C(5)C 116.6(0.5) O(22)C–N(21)C–O(23)C 124.4(0.5) C(2)C–C(1)C–C(8)C 176.4(0.6) O(23)C–N(21)C–C(5)C 119.0(0.5) C(1)C–C(2)C–C(7)C 121.7(0.5) C(1)C–C(2)C–C(3)C 118.8(0.5) N(15)C–C(3)C–C(2)C 117.2(0.5) C(3)C–C(2)C–C(7)C 119.6(0.5) C(2)C–C(3)C–C(4)C 120.0(0.5) N(15)C–C(3)C–C(4)C 122.8(0.5) N(21)C–C(5)C–C(4)C 115.4(0.5) C(3)C–C(4)C–C(5)C 117.3(0.5) C(4)C–C(5)C–C(6)C 124.6(0.5) N(21)C–C(5)C–C(6)C 120.0(0.5) C(2)C–C(7)C–C(6)C 120.7(0.5) C(5)C–C(6)C–C(7)C 117.8(0.5) C(8)C–C(9)C–C(10)C 122.3(0.5) C(1)C–C(8)C–C(9)C 177.9(0.6) C(10)C–C(9)C–C(14)C 119.2(0.6) C(8)C–C(9)C–C(14)C 118.6(0.5) C(10)C–C(11)C–C(12)C 120.0(0.6) C(9)C–C(10)C–C(11)C 121.3(0.6) O(24)C–C(12)C–C(13)C 124.4(0.5) O(24)C–C(12)C–C(11)C 115.6(0.5) C(12)C–C(13)C–C(14)C 119.9(0.6) C(11)C–C(12)C–C(13)C 120.0(0.6) O(17)C–C(16)C–C(20)C 124.7(0.5) C(9)C–C(14)C–C(13)C 119.5(0.6) O(20)C–C(16)C–N(15)C 127.8(0.5) O(17)C–C(16)C–N(15)C 107.6(0.5) O(24)C–C(25)C–C(26)C 107.8(0.5) O(17)C–C(18)C–C(19)C 105.1(0.5) C(25)C–C(26)C–C(31)C 117.2(0.6) C(25)C–C(26)C–C(27)C 121.3(0.6) C(26)C–C(27)C–C(28)C 119.5(0.6) C(27)C–C(26)C–C(31)C 121.5(0.6) C(28)C–C(29)C–C(30)C 118.6(0.6) C(27)C–C(28)C–C(29)C 121.4(0.7) C(26)C–C(31)C–C(30)C 116.7(0.6) C(29)C–C(30)C–C(31)C 122.2(0.6) aE.s.d.s shown in parentheses.crystals. The largest parallel surfaces were found to be (10-1) and that of the X-axis and the optic axis was 39° at 532 nm. Assuming that the directions do not depend on the wavelength, by X-ray diraction. Since the crystal is monoclinic, the b axis and the Y-axis coincide and are found to align in the direction nx and nz at three wavelengths were calculated from the measured refractive indices in the direction of surface normal.Fig. 3 of the pillar. Although the optical Y -axis was in the surface of the measured crystal, X- and Z-axes were not in the surface. shows these values and fitted curves of the Sellmeier equation. Fig. 4(a)–(c) shows a molecule from the a-form crystal and The angle between the X-axis and the surface normal was 12° 708 J.Mater. Chem., 1997, 7(5), 705–711the packing of the molecules in the crystal lattice viewed along the b and c axes, respectively. The angle between the two benzene rings bridged by the acetylene moiety is 10.6°. As is clearly seen in Fig. 4(c), the molecules in a unit cell are aligned in two directions in the ab plane.The angle between them is 88.0° and the angle between a molecule and the a-axis is 44.0°. The distance between O(20) and N(15) of the neighbouring molecule one unit cell along the b-axis is 2.933(7) A° . The hydrogen bond N(15)–H(N15),O(20) is formed between two neighbouring molecules, and is therefore expected to be one of the factors that destroys the centrosymmetric structure.In the b-form crystal, there are three crystallographically independent molecules, which are represented here by molecules A, B, and C. The structures of these molecules are shown in Fig. 5(a). For all three molecules, the two benzene rings linked by the triple-bonded carbons are in approximately the same plane, which is almost normal to the ac plane. However, the benzene ring in the alkoxy group is tilted 20° Fig. 3 Refractive indices of 1 in the a-form crystal. Closed circles are from the plane of other benzene rings in molecule A, 74° the measured value and the open circles are the calculated values. The in the molecule B and 62° in the molecule C. In Fig. 5(b), solid curves denote the value fitted Sellmeier equations. which is the packing diagram of one unit cell viewed along Fig. 4 (a) Molecular diagram of the a-form crystal. (b) The molecular packing diagram of the a-form crystal viewed along the b-axis. (c) The molecular packing diagram of the a-form crystal viewed along the c-axis. J. Mater. Chem., 1997, 7(5), 705–711 709Fig. 5 (a) Molecular diagram of the b-form crystal. (b) The molecular packing diagram of the b-form crystal viewed along the b-axis.(c) The molecular packing diagram of the b-form crystal viewed along the c-axis. the b-axis, the stacking of the molecules is clear. The view than ethyl acetate (e=6.02). For crystallization of 1, the strong polar solvent aects the formation of a non-centrosymmetric along the c-axis is also shown [Fig. 5(c)]. The hydrogen atom on N(15) and the neighbouring O(20) are too far apart to crystal.13 Considering that the crystal from a melt is always in the c-form and finally turns into the b-form, the b-form crystal form hydrogen bonds.We therefore concluded that the hydrogen bond, which plays an important role in determining the must be more stable than the others. However, by using a strong polar solvent during the crystal growth it was possible crystal structure, does not exist in the b-form crystal. Molecule 1 is designed to have a large charge transfer to make a less stable crystal. The nonlinear optical properties of crystals of 1 were meas- between the benzyloxy group and the nitro group through the benzene rings and the triple bond.This charge transfer leads ured using the SHEW technique.The SHEW power from the to a large hyperpolarizability. On the other hand, such mol- b-form crystal was negligibly small, while on the other hand, ecules tend to align antiparallel to each other because of the the a-form crystal showed high SH activity. The SHEW power large ground-state dipole moments. If the molecules were too from the a-form crystal was about ten times that from m-NA, close to each other they would repulse each other and align which means that the nonlinear optical coecient of 1 should in opposite directions.The carbamate group was introduced be more than three times that of m-NA. Thus compound 1 is for steric reasons, so that the molecules do not align and the expected to be a high SH active material. dipole–dipole interaction is decreased.Additionally, the hydrogen bond between the nitrogen in the amide group and the oxygen in the same group of the next molecule causes the non- Conclusion centrosymmetric structure, as shown in the a-form crystal We have synthesized ethyl 2-(4-benzyloxyphenylethynyl)-5- structure. The carbamate group aects not only the nonlinnitrobenzene- 1-carbamate and studied its polymorphism and earity of a molecule but also the crystal structure, as this group crystal structures.Three types of crystal were found which makes the molecule bulky and the nonlinear eciency in the were interchangeable by heating and cooling the sample. One unit volume is decreased. Comparing the recrystallization of of the crystals, crystallized from a polar solvent, shows non- the a- and b-forms from solution, the relative permittivities of centrosymmetry, which is induced by the hydrogen bonds the solvents used are dierent. 2-Ethoxyethanol, from which a-form crystallizes, has a larger relative permittivity (e=29.6) between molecules. It is expected to have a large nonlinear 710 J. Mater. Chem., 1997, 7(5), 705–711J. Phys. Chem., 1992, 96, 6232; (c) T.Kurihara, H. Tabei and optical coecient which is more than three times that of T. Kaino, J. Chem. Soc., Chem. Commun., 1987, 959; meta-nitroaniline. (d) A. E. Stiegman, E. Graham, K. J. Perry, L. R. Khundkar, L.- T. Cheng and J. W. Perry, J. Am. Chem. Soc., 1991, 113, 7658. We gratefully thank Dr Takashi Kondo and Professor Ryoichi 7 (a) T. M. Robertson and I. Woodward, Proc.R. Soc. L ondon, Ser. Ito of Tokyo University for their help with refractive index A, 1937, 162, 436; (b) S. D. Samarskaya, R. M. Myasnikova and A. I. Kitiagorodkii,Kristallografiya, 1968, 13, 616; (c) A. A. Espiritu measurements. We thank Dr Takao Tomono of Fuji Xerox and J. G. White, Z. Kristallogr., 1977, 147, 177; (d) A. A. Espiritu Co., Ltd., for his useful comments and advice. and J. G. White, Acta Crystallogr., 1977, B33, 3899; (e) E. M. Graham, V. M. Miskowaski, J. W. Perry, D. R. Coulter, A. E. Stiegman, W. P. Shaefer and R. E. Marsho, J. Am. Chem. References Soc., 1989, 111, 8771. 8 M. Kiguchi, M. Kato, N. Kumegawa and Y. Taniguchi, J. Appl. 1 Non-linear optical properties of organic and polymeric materials, ed. Phys., 1994, 75, 4332. D. J. Williams, Am. Chem. Soc. Symp. Ser. 233, American 9 M. Kiguchi, M. Kato, N. Kumegawa, M. Okunaka and Chemical Society,Washington DC, 1983. Y. Taniguchi, Nonlinear Optics, 1995, 9, 223. 2 Nonlinear optical properties of organic molecules and crystals, ed. 10 G. W. Gray and B. Jones, J. Chem. Soc., 1954, 1467. D. S. Chemla and J. Zyss, Academic Press, Orlando, 1987. 11 A. D. Allen and C. D. Cook, Can. J. Chem., 1963, 41, 1084. 3 B. F. Levine and C. G. Betha, J. Chem. Phys., 1977, 66, 1070. 12 J. J. Blanksma, Recl. T rav. Chim. Pays-Bas, 1946, 65, 329. 4 J. L. Oudar and D. S. Chemla, J. Chem. Phys., 1977, 66, 2664. 13 H. Tabei, T. Kurihara and T. Kaino, Appl. Phys. L ett., 1987, 50, 5 S. J. Lalama and A. F. Garito, Phys. Rev. A, 1979, 20, 1179. 1855. 6 (a) M. Nakano, K. Yamaguchi and T. Fueno, Springer Proceedings in Physics, vol. 36, Nonlinear Optics of Organics and Semiconductors, 1989, 98; (b) N. Matsuzawa and D. A. Dixon, Paper 6/07858C; Received 19th November, 1996 J. Mater. Chem., 1997, 7(5), 705–711 711

ISSN:0959-9428    

DOI:10.1039/a607858c

出版商:RSC    

年代:1997

数据来源: RSC

摘要:

A novel class of phenol–pyridine co-crystals for second harmonic generation Kin-Shan Huang,a† Doyle Britton,a the late Margaret C. Ettera and Stephen R. Byrn*b aDepartment of Chemistry University of Minnesota Minneapolis,MN 55455 USA bDepartment ofMedicinal Chemistry and Pharmacognosy Purdue University IN 47907 USA To test the approach of combining both ionic and hydrogen-bonding interactions for the design of non-linear optical (NLO) materials a number of phenol–pyridine co-crystals have been synthesized and their NLO properties investigated. The co-crystals are characterized by second harmonic generation measurements as well as the more conventional methods of melting point measurements infrared and nuclear magnetic resonance spectroscopy. To investigate whether the phenol co-crystals are organic salts the 2-methoxy-4-nitrophenol–4-(dimethylamino)pyridine (251) co-crystal 6 and the 2-methoxy-4-nitrophenol–4-pyrrolidinylpyridine –water (15151) co-crystal 8 are further characterized by X-ray single-crystal diraction.Crystal structure analyses reveal that both 6 and 8 are ionic co-crystals (or organic salts) composed of a phenoxide anion a pyridinium cation and a neutral molecule. In the two co-crystals the phenoxide pyridinium and neutral molecules are held together by ionic attractions as well as hydrogen-bonding interactions. Both 6 and 8 crystallize in non-centrosymmetric structures [Pna21 (orthorhombic) a=6.880(4) b=38.40(1) c=8.454(3) A° Z=4 Dc=1.369 g cm-3 and R=0.051 for 6 and Cc (monoclinic) a=7.302(3) b=23.518(2) c=9.940(1) A° b=107.12(2)° Z=4 Dc=1.365 g cm-3 and R=0.036 for 8].In addition to X-ray structure determination it is possible to predict whether phenol–pyridine co-crystals are organic salts based on the DpKa [pKa(pyridine)-pKa(phenol)] and stoichiometric ratio of the co-crystals. Preliminary results suggest that this type of co-crystals particularly for the ionic co-crystals may have a higher chance of forming non-centrosymmetric structures than the normal achiral organic compounds. The design and search for non-linear optical (NLO) materials the tendency of the organic dipolar molecules to form antiparallel pairs as discussed earlier this eect may be intensified by has been of interest in recent years because of their potential applications to laser devices telecommunications and optical combining the ionic interactions with hydrogen-bonding interactions.One way to include both ionic and hydrogen-bonding data storage.1 Second harmonic generation (SHG) the secondorder eect of NLO properties of crystalline materials depends interactions in a system is to form ionic co-crystals from proton transfer between an organic acid and organic base as both on the magnitude of the molecular hyperpolarizability (b) (microscopic non-linearity) and on the orientation of the illustrated in eqn. (1) where A–H is an organic acid B is an organic base and A-,H–B+ is the resulting ionic co-crystal molecules in the crystal lattice.2 Organic molecules in general are potentially more attractive and versatile than inorganic (or organic salt) which involves in both ionic attractions and hydrogen-bonding interactions. compounds for NLO materials because of their large b-values fast response time high resistance to optical damage and the almost unlimited possibilities of designing molecules suitable A–H+B CCCCCA co-crystallization A-,H–B+ (1) for SHG.3 Although many organic molecules such as p- In this study this approach was tested by complexing a nitroaniline that contain both electron-donating and electron- number of cyanophenol–pyridine and nitrophenol–pyridine accepting groups linked via aromatic rings have inherently pairs to form co-crystals.Since cyanophenols and nitrophenols large b-values they do not exhibit SHG in the solid state since are both good hydrogen-bond donors and strong organic acids these molecules pack into centrosymmetric structures for which and since pyridines especially disubstituted 4-aminopyridines the macroscopic non-linearity x(2) is necessarily zero.4 Since are both good hydrogen-bond acceptors and strong organic organic molecules with large b-values are usually accompanied bases the phenol–pyridine pairs are expected to form ionic by high molecular dipole moments they tend to form pairs co-crystals with proton transfer between the hydroxy hydrogen aligned in an anti-parallel fashion and thus tend to crystallize of the phenols and the pyridine nitrogen of the pyridines.In in centrosymmetric space groups.5 One strategy employed to addition cyanophenols and nitrophenols are well known for encourage non-centrosymmetric crystallization of these mol- their large intrinsic molecular hyperpolarizability so their co- ecules is to increase their asymmetry by introducing a substitu- crystals may exhibit significant SHG if they crystallize in non- ent in the meta position of the aromatic rings.6 An alternative centrosymmetric space groups.Interestingly a similar class of approach first proposed by Meridith is the use of ionic ionic complexes formed by inorganic acids (e.g. phosphoric interactions to override the deleterious dipolar interactions acid) and organic bases (e.g. amines) has been reported and drive the formation of salts in non-centrosymmetric crystal recently.12 In this paper besides the preparation and charac- structures.7 Marder and co-workers have further explored the terization of the phenol–pyridine co-crystals whether the co- methodology to investigate a number of 4-N-methylstilbazol- crystals are organic salts was also investigated.Finally the ium salts and found that the ionic compounds exhibit a higher chance for this type of co-crystals to form non-centrosymmetric tendency to crystallize in non-centrosymmetric crystal structures was evaluated. structures than do neutral dipolar compounds.8 Hydrogen bonds are strong directive and predictable noncovalent interactions. The importance of hydrogen bonding in Experimental the crystal packing of organic compounds has long been General methods recognized.9–11 Since ionic interactions are useful to counteract A Nikon optical microscope equipped with polarizing filters was used for microscopic investigation. Melting points were † Current address Pharmaceutical Research Warner-Lambert Company 170 Tabor Road Morris Plains NJ 07950 USA. determined on a Fisher-Johns apparatus and are uncorrected.J. Mater. Chem. 1997 7(5) 713–720 713 Infrared (IR) spectra were recorded on a Nicolet 5DXB FTIR 4-Nitrophenol–trans-1,2-bis( 4-pyridyl )ethylene (251) co-crystal 3 spectrometer from 4000 to 450 cm-1 using Nujol mulls. 1H nuclear magnetic resonance (NMR) spectra were recorded on Methods A B and C; o-white needles (method A ethyl an IBM NR200AF spectrometer and chemical shifts are acetate); mp 166–167 °C; n/cm-1 2566 (br) 1820 (br) 1607 reported in parts per million (d) relative to tetramethylsilane 1590 1496 1466 1336 1299 1255 1115 1604 1011 847 827 (Me4Si). Measurements of powder SHG eciencies were per- and 753; 1H NMR (CDCl3) d 7.03 (dd 4H) 7.53 (s 2H) 7.59 formed on a modified apparatus using the Kurtz powder (dd 4H) 8.16 (dd 4H) and 8.60 (dd 4H); SHG 0.technique13 with the fundamental wavelength (1074 nm) of a Q-switched Nd:YAG laser. Except for 4-(methylamino)pyrid- 4-Nitrophenol–trans-1,2-bis( 4-pyridyl )ethane (151) co-crystal ine which was synthesized according to Wibaut and 4 Broekman’s method,14 all nitrophenols cyanophenols and pyridines were purchased from Aldrich Chemical Company Methods A B and C; o-white blocks (method A ethyl (Milwaukee WI) and recrystallized from appropriate solvents. acetate); mp 159–160 °C; n/cm-1 2558 (br) 1813 (br) 1613 Spectroscopic grade solvents were obtained from Fischer 1587 1517 1494 1461 1379 1335 1290 1268 1249 1227 Scientific (Pittsburgh PA) and used without further purifi- 1108 1011 844 and 833; 1H NMR (CDCl3) d 3.01 (s 4H) cation. Deuteriated solvents were also obtained from Aldrich 6.90 (dd 2H) 7.17 (dd 4H) 8.13 (dd 2H) and 8.51 (dd 4H); Chemical Company.SHG 0. General methods for the preparation of phenol–pyridine 4-Nitrophenol–4-phenylpyridine (151) co-crystal 5 co-crystals Methods A B and C; colourless plates (method A ethyl The co-crystals were prepared from various phenol–pyridine acetate); mp 88–90°C; n/cm-1 2566 (br) 1805 (br) 1595 1519 pairs by solution and/or solid-state methods as described 1501 1468 1335 1307 1275 1114 1108 1011 833 760 and below. 617; 1H NMR (CDCl3) d 6.93 (dd 2H) 7.51 (dd 2H) 7.67 (m 4H) 8.14 (dd 2H) and 8.66 (dd 2H); SHG 0. Method A (slow evaporation). Stoichiometric amounts of the phenols and pyridines were dissolved in a mutually miscible 2-Methoxy-4-nitrophenol–4-(dimethylamino)pyridine (251) solvent system and allowed to recrystallize at room tempera- co-crystal 6 ture.Whenever possible the resulting crystals were removed Methods A B and C; long yellow needles (method B ethyl from solution as soon as they were formed before the solvent acetate–hexane); mp 99–101°C; n/cm-1 2651 (br) 2031 1926 had completely evaporated. Alternatively co-crystals were 1644 1553 1483 1335 1314 1265 1223 1089 1026 857 829 obtained after all of the solvent had evaporated. 808 787 and 745; 1H NMR (CDCl3) d 3.09 (s 6H) 3.61 (s 6H) 6.58 (dd 2H) 6.87 (d 2H) 7.70 (d 2H) 7.81 (dd 2H) Method B (vapour diusion). In a 25 ml beaker stoichiometric 8.11 (dd 2H) and 8.66 (dd 2H); SHG 2×urea. amounts of the phenols and pyridines were dissolved in a mutually miscible solvent system. The beaker was placed in a 2-Methoxy-4-nitrophenol–4-pyrrolidinylpyridine (251) 200 ml jar and a second solvent was introduced into the jar as co-crystal 7 a precipitant.The jar was sealed and left to stand undisturbed. As the precipitant diused into the solution crystals formed Methods B and C; yellow needles (method B ethyl acetate– in the solution. hexane); mp 104–111°C; n/cm-1 3114 2024 1926 1652 1616 1553 1518 1490 1462 1328 1300 1258 1209 1096 1026 864 Method C (solid-state grinding). Stoichiometric amounts of 829 808 702 and 639; 1H NMR (CDCl3 ) d 2.08 (quintet 4H) the phenols and pyridines were ground together in a Wig-L- 3.38 (t 4H) 3.91 (s 6H) 6.46 (d 2H) 6.87 (d 2H) 7.71 Bug dental amalgamator for 10 min. Conversion of the starting (d 2H) 7.81 (dd 2H) and 8.08 (d 2H); SHG 0. materials to co-crystals in the solid state often approached 100%.2-Methoxy-4-nitrophenol–4-pyrrolidinylpyridine–water (15151) Co-crystals prepared by all methods were characterized by co-crystal 8 optical microscopy melting point measurements IR and 1H Method A; yellow needles (ethyl acetate); mp 68–77 °C; n/cm-1 NMR spectrometry and powder SHG measurements. In all 3585 3332 3241 2651 (br) 1644 1602 1560 1511 1462 1420 cases where two or more methods were used for one system 1350 1277 1250 1215 1198 1183 1125 1092 1026 862 826 the same solid-state phases were obtained from the dierent 803 and 636; 1H NMR [(CD3)2CO] d 2.04 (quintet 4H) 3.40 methods. The methods that had been used successfully to (t 4H) 3.95 (s 3H) 5.06 (s 3H) 6.49 (dd 2H) 6.94 (d 1H) prepared co-crystals are listed below. The crystal morphology 7.75–7.85 (m 2H) and 8.12 (d 2H); SHG 2×urea.of the co-crystals is described and the method used for growing single crystals is indicated in parentheses. 2-Methoxy-5-nitrophenol–4-(dimethylamino)pyridine (151) co-crystal 9 4-Nitrophenol–4-(dimethylamino)pyridine (251) co-crystal 1 Method B; yellow plates (ethyl acetate–hexane); mp 74–76°C; Methods A B and C; yellow needles (method B ethyl acetate– n/cm-1 1652 1616 1595 1511 1448 1349 1293 1272 1223 hexane); mp 128–130 °C; n/cm-1 1644 1578 1557 1490 1467 1082 1019 1005 815 and 714; 1H NMR (CDCl3) d 3.02 (s 1378 1323 1271 1254 1213 1173 1104 994 850 824 755 6H) 3.96 (s 3H) 6.52 (dd 1H) 6.86 (d 1H) 7.74 (d 1H) 7.76 and 643; 1H NMR (CDCl3) d 3.13 (s 6H) 6.62 (d 2H) 6.89 (d 1H) 8.19 (dd 2H) and 9.25 (br s 1H); SHG 0. (d 4H) 8.13 (dd 6H) and 8.71 (brs 2H); SHG 2×urea.4-Nitrophenol–4-pyrrolidinylpyridine (151) co-crystal 2 2-Methoxy-5-nitrophenol–4-(dimethylamino)pyridine (251) co-crystal 10 Methods A and C; yellow plates (method A ethyl acetate); mp 89–91°C; n/cm-1 2369 (br) 1912 (br) 1627 1608 1587 1517 Method B and C; deep yellow plates (method B ethyl acetate– hexane); mp 57–63°C; n/cm-1 3121 1652 1595 1560 1525 1499 1459 1332 1295 1219 1110 1009 857 and 808; 1H NMR (CDCl3) d 2.07 (quintet 4H) 3.37 (t 4H) 6.45 (dd 2H) 1497 1469 1335 1272 1251 1216 1082 1019 815 and 745; 1H NMR (CDCl3) d 3.05 (s 6H) 3.96 (s 6H) 5.62 (s 4H) 6.85 (dd 2H) and 8.09 (m 4H); SHG 0. 714 J. Mater. Chem. 1997 7(5) 713–720 6.54 (dd 2H) 6.86 (d 2H) 7.75–7.80 (m 4H) and 8.16 (dd (t 4H) 3.33 (t 4H) 3.86 (s 3H) 6.41 (dd 2H) 6.89 (d 1H) 7.03 (d 1H) 7.14 (dd 1H) and 8.12 (dd 2H); SHG 0.2H); SHG 0. 2,4-Dinitrophenol–4-(dimethylamino)pyridine (151) co-crystal 2-Methoxy-5-nitrophenol–4-pyrrolidinylpyridine (151) 19 co-crystal 11 Methods A and C; yellow needles (method A 151 MeOH– Methods A B and C; orange needles (method B ethyl acetate– ethyl acetate); mp 169–171°C; n/cm-1 2052 1954 1644 1616 hexane); mp 89–91°C; n/cm-1 1625 1609 1532 1511 1467 1588 1547 1534 1464 1441 1373 1318 1267 1247 1212 1450 1335 1286 1237 1082 1012 871 808 and 745; 1H NMR 1135 and 825; 1H NMR (CDCl3) d 3.22 (s 6H) 6.69 (d 2H) (CDCl3) d 2.03 (quintet 4H) 3.32 (t 4H) 3.95 (s 3H) 6.40 6.79 (d 1H) 8.07 (dd 1H) 8.23 (d 2H) and 9.00 (d 1H); (dd 2H) 6.84 (dd 1H) 7.71–7.77 (m 2H) 8.15 (dd 2H) and SHG 0. 10.08 (br s 1H); SHG 0. 2,4-Dinitrophenol–4-pyrrolidinylpyridine (151) co-crystal 20 4-Nitrophenol–2-methoxy-4-nitrophenol–4- (dimethylamino)pyridine (15151) co-crystal 12 Methods A and C; yellow needles (method A 151 MeOH– ethyl acetate); mp 177–179 °C; n/cm-1 2644 (br) 2031 1933 Method B; yellow prisms (ethyl acetate–hexane); mp 96–99 °C; 1652 1595 1553 1532 1469 1370 1321 1250 1209 1131 and n/cm-1 1644 1609 1588 1560 1504 1490 1469 1335 1251 815; 1H NMR (CDCl3) d 2.15 (quintet 4H) 3.48 (quintet 4H) 1209 1096 1026 875 815 and 787; 1H NMR (CDCl3) d 3.11 6.56 (d 2H) 6.71 (d 1H) 8.02 (dd 1H) 8.16 (d 2H) and 8.96 (s 6H) 3.79 (s 3H) 6.61 (d 2H) 6.85–6.92 (m 3H) 7.72 (d 1H); SHG 0.(d 1H) 7.83 (dd 1H) 8.06 (d 2H) and 8.09 (dd 2H); SHG 0. 4-Nitrophenol–4-(methylamino)pyridine (151) co-crystal 21 4-Nitrophenol–2-methoxy-4-nitrophenol–4-pyrrolidinylpyridine (15151) co-crystal 13 Method A; yellow plates (ethyl acetate); mp 133–137°C; n/cm-1 3374 1644 1581 1546 1469 1377 1335 1002 and 1103; 1H Method B; yellow needles (ethyl acetate–hexane); mp NMR (CDCl3) d 2.90 (s 3H) 6.68 (dd 2H) 6.94 (tt 4H) and 112–114 °C; n/cm-1 1652 1606 1588 1560 1514 1501 1459 8.08–8.19 (m 6H); SHG 0.1335 1257 1201 1091 861 and 829; 1H NMR (CDCl3) d 2.09 (quintet 4H) 3.39 (t 4H) 3.92 (s 3H) 6.48 (d 2H) 6.83 2-Methoxy-4-nitrophenol–4-(methylamino)pyridine (151) (d 1H) 6.88 (tt 2H) 7.71 (d 1H) 7.83 (dd 1H) and 8.04– co-crystal 22 8.09 (m 4H); SHG 0. Methods A and C; deep yellow needles (method A ethyl 2-Methoxy-4-nitrophenol–2-methoxy-5-nitrophenol–4- acetate); mp 115–118°C; n/cm-1 3227 3107 1652 1622 1574 pyrrolidinylpyridine (15151) co-crystal 14 1546 1518 1462 1356 1284 1260 1245 1218 1191 1096 1033 822 and 639; 1H NMR (CDCl3) d 2.88 (s 3H) 3.89 (s Method B; yellow needles (ethyl acetate–hexane); mp 93–95 °C; 3H) 5.10 (br s 1H) 6.42 (dd 2H) 6.77 (d 1H) 7.70 (d 1H) n/cm-1 3121 2038 1940 1652 1616 1586 1573 1556 1504 7.80 (dd 1H) and 8.09 (dd 2H); SHG 0.1461 1435 1354 1332 1272 1255 1216 1096 1023 860 830 and 804; 1H NMR (CDCl3 ) d 2.06 (quintet 4H) 3.36 (t 4H) Crystal structure determinations for 6 and 8 3.90 (s 3H) 3.95 (s 3H) 6.45 (dd 2H) 6.83 (s 1H) 6.87 (t 1H) 7.70–7.84 (m 4H) and 8.11 (dd 2H); SHG 0. Single crystals of 6 and 8 were obtained as yellow plates and yellow prisms respectively and identified by IR and 1H NMR 4-Cyanophenol–4-(dimethylamino)pyridine (151) co-crystal 15 spectrometry.The crystal used for the diraction studies was 0.50×0.30×0.25 mm for 6 and 0.60×0.50×0.30 mm for 8. Methods A B and C; white plates (method B ethyl acetate– Crystallographic data for both structures are given in Table 1. hexane); mp 102–104 °C; n/cm-1 2214 1609 1581 1532 1511 Preliminary examinations and data collections were 1461 1384 1293 1223 1116 1005 850 and 808; 1H NMR performed with an Enraf-Nonius CAD-4 single-crystal X-ray (CDCl3) d 3.04 (s 6H) 6.53 (dd 2H) 6.84 (d 2H) 7.40 (d 2H) and 8.11 (dd 2H); SHG 0. Table 1 Crystallographic data for 6 and 8 4-Cyanophenol–4-pyrrolidinylpyridine (151) co-crystal 16 6 8 Methods A B and C; white blocks (method B ethyl acetate– hexane); mp 83–85°C; n/cm-1 2214 1609 1574 1532 1511 chemical formula C21H24N4O8 C16H20N3O5 1462 1405 1293 1216 1173 1005 857 and 808; 1H NMR formula weight 460.44 335.35 crystal system orthorhombic monoclinic (CDCl3) d 2.04 (quintet 4H) 3.32 (t 4H) 6.42 (dd 2H) 6.86 space group Pna21 Cc (tt 2H) 7.43 (tt 2H) and 8.08 (dd 2H); SHG 0.a/A° 6.880(4) 7.302(3) b/A° 38.40(1) 23.518(2) 4-Hydroxy-3-methoxybenzonitrile–4-(dimethylamino)pyridine c/A° 8.454(3) 9.940(1) (151) co-crystal 17 b/degrees — 107.12(2) Z 4 4 Methods A B and C; white thin plates (method B ethyl V /A° 3 2233(4) 1631(1) acetate–hexane); mp 68–70°C; n/cm-1 2207 1616 1600 1543 Dc/g cm-3 1.369 1.365 1517 1467 1446 1377 1272 1216 1005 and 808; 1H NMR m(Mo–Ka)/cm-1 1.00 0.96 F(000) 968 708 (CDCl3) d 2.97 (s 6H) 3.80 (s 3H) 6.47 (dd 2H) 6.84 number of reflections collected 4131 3262 (d 1H) 6.98 (d 1H) 7.08 (dd 1H) and 8.11 (dd 2H); SHG 0.Rint 0.031 0.015 reflections with I>2.0s(I) 1295 1416 4-Hydroxy-3-methoxybenzonitrile–4-pyrrolidinylpyridine (151) number of variables 294 224 co-crystal 18 R; wR 0.051; 0.068 0.036; 0.045 goodness of fit 2.18 1.36 Methods A B and C; white needles (method B ethyl acetate– maximum shift/esda 0.07 0.07 hexane); mp 119–121 °C; n/cm-1 2221 1609 1525 1462 1405 1279 1159 1131 1012 and 815; 1H NMR (CDCl3 ) d 2.04 aesd=estimated standard deviation. J. Mater. Chem. 1997 7(5) 713–720 715 diractometer with Mo-Ka radiation (l=0.71069 A° ). The noncentrosymmetric space groups of 6 (Pna21) and 8 (Cc) were determined based on the systematic absence of 0kl k+l2n and h0l h2n (for 6) and hkl h+k2n and h0l l2n (for 8) and were confirmed by the positive powder SHG test. Data collection on the diractometer was performed at 24°C using the v scan technique to a maximum 2h of 48.0° for 6 and 51.9° for 8.All data were corrected for Lorentz and polarization eects. An empirical absorption correction using the DIFABS program15 was applied to 6 only which resulted in transmission factors ranging from 0.73 to 1.25. The crystal structures were solved by direct methods.16,17 All non-hydrogen atoms were refined anisotropically. Except for the NH and OH protons which were refined isotropically the other hydrogen atoms were included in the structure factor calculations and placed in idealized position (dC—aH=0.95 A° ) with assigned isotropic thermal parameters (B=1.2 B of bonded atoms). Full-matrix least-square refinements were based on 1295 observed reflections [I>2.0s(I)] with 294 variable parameters for 6 and 1416 observed reflections [I>2.0s(I)] with 224 variable parameters for 8.Scattering factors for neutral atoms and Df ¾ and Df were taken from ref. 18. All calculations were performed using a TEXSAN crystallographic software package.19 Atomic coordinates thermal parameters and bond lengths and angles have been deposited at the Cambridge Fig. 1 The crystal structure and atomic numbering of 6. Hydrogen Crystallographic Data Centre (CCDC). See Information for bonds are indicated by dashed lines. Authors J. Mater. Chem. 1997 Issue 1. Any request to the CCDC for this material should quote the full literature citation is composed of one 2-methoxy-4-nitrophenoxide anion one 4- and the reference number 1145/19. (dimethylamino)pyridinium cation and one neutral 2-methoxy- 4-nitrophenol molecule.The atomic numbering scheme of 6 is shown in Fig. 1. Selected bond distances bond angles and Results and Discussion torsion angles of the co-crystal are listed in Table 2. In this structure the phenoxide and pyridinium ions and Twenty-two phenol co-crystals were prepared by complexing neutral phenol molecules are all planar. The phenoxide and a variety of phenol–pyridine pairs includingtwo cyanophenols pyridinium ions are nearly co-planar to each other with a four nitrophenols and five pyridines. Since a phenol–pyridine dihedral angle of 5.2° between them. The neutral phenol co-crystal is not a simple mixture of the phenol and pyridine molecule however is nearly perpendicular to both the phen- molecules but a new solid-state material the co-crystal usually oxide and pyridinium ions.The dihedral angles between the has dierent physical properties such as crystal morphology aromatic plane of the neutral phenol and those of the phen- melting point IR spectrum and X-ray diraction pattern from oxide and pyridinium ions are 73.9 and 72.5° respectively. its starting materials. All the co-crystals prepared were charac- Since the hydroxy group of the neutral phenol is not oriented terized by optical microscope melting point measurements toward the ortho-substituted methoxy group this hydroxy solid-state IR and solution 1H NMR spectrometry X-ray group does not form an intramolecular hydrogen bond to the powder diraction (data not shown) and powder SHG methoxy oxygen atom. The occurrence of proton transfer is measurements.evident by the short N(2)MH(2N) bond length of 1.092(7) A° Solid-state IR was particularly useful for characterizing the (Table 2) although this bond length is slightly longer than the phenol co-crystals. Since the OMH stretching bands of the average neutral NMH bond length (0.938 A° ).21 As a conse- hydroxy group in the host phenols are very sensitive to quence of the proton transfer the bond lengths of hydrogen-bonding environment,20 formation of the phenol co- C(4B)MO(3B) [1.317(9) A° ] and C(1B)MN(1B) [1.389(9) A° ] crystals can be monitored by comparing the OMH stretching show a significant decrease compared to those of bands in the phenols with those in their co-crystals. The strong C(4A)MO(3A) [1.360(8) A° ] and C(1A)MN(1A) [1.47(1) A° ] OMH stretching band (3367 cm-1) of 2-methoxy-4-nitrophein the neutral phenol (Table 2).These results suggest that the nol disappeared and a weak broad band near 2650 cm-1 was contribution of the quinoid form II shown in Fig. 2 increases observed in 6. The broad band could be either from the strong significantly in the structure of the phenoxide ion. hydrogen-bonded OMH stretching of the neutral phenol or Because of the contribution of the quinoid form the bond the +NMH stretching of the protonated pyridinium anion or lengths of C(2B)MC(3B) [1.35(1) A° ] and C(5B)MC(6B) the combination of the two. Similarly the shift of the OMH [1.37(1) A° ] are noticeably shorter than the other CMC bond stretching bands to lower wavelengths was also observed in lengths in the quinoidphenyl ring C(1B)MC(2B) [1.432(9) A° ] other phenol co-crystals.The formation of phenol co-crystals C(3B)MC(4B) [1.40(1) A° ] C(4B)MC(5B) [1.39(1) A° ] and can also be monitored by 1H NMR spectroscopy if single C(1B)MC(6B) [1.41(1) A° ]. In addition to the ionic attraction crystals are available. Based on 1H NMR studies the stoichiometric ratios of the co-crystals were determined and are included in the experimental section. To investigate whether the co-crystals formed are ionic salts co-crystals 6 and 8 were further characterized by X-ray structure determination. Crystal structure analysis of 6 This co-crystal is an organic salt because proton transfer has Fig. 2 occurred between the phenol and pyridine. This organic salt 716 J. Mater. Chem. 1997 7(5) 713–720 Table 2 Selected bond distances (A° ) bond angles (degrees) and torsion angles (degrees) in 6 C(4A)–O(3A) 1.360(8) C(4B)–O(3B) 1.317(9) C(1A)–N(1A) 1.47(1) C(1B)–N(1B) 1.389(9) C(5A)–O(4A) 1.392(9) C(5B)–O(4B) 1.369(8) C(7A)–O(4A) 1.409(9) C(7B)–O(4B) 1.462(8) C(1A)–C(2A) 1.34(1) C(1B)–C(2B) 1.432(9) C(1A)–C(6A) 1.40(1) C(1B)–C(6B) 1.41(1) C(2A)–C(3A) 1.35(1) C(2B)–C(3B) 1.35(1) C(3A)–C(4A) 1.40(1) C(3B)–C(4B) 1.40(1) C(4A)–C(5A) 1.370(9) C(4B)–C(5B) 1.39(1) C(5A)–C(6A) 1.34(1) C(5B)–C(6B) 1.37(1) C(8)–C(9) 1.36(1) C(11)–C(12) 1.34(1) C(8)–C(10) 1.36(1) C(10)–C(12) 1.42(1) C(9)–N(2) 1.32(1) C(11)–N(2) 1.33(1) C(13)–N(3) 1.46(1) C(14)–N(3) 1.46(1) C(10)–N(3) 1.36(1) N(2)–H(2N) 1.092(7) O(1A)–N(1A)–O(2A) 121.8(8) O(1B)–N(1B)–O(2B) 120.0(7) C(5A)–O(4A)–C(7A) 117.6(6) C(5B)–O(4B)–C(7B) 117.1(5) C(3A)–C(4A)–O(3A) 122.4(7) C(3B)–C(4B)–O(3B) 120.4(7) C(5A)–C(4A)–O(3A) 119.5(7) C(5B)–C(4B)–O(3B) 121.0(7) C(10)–N(3)–C(13) 121.2(6) C(10)–N(3)–C(14) 120.3(7) C(13)–N(3)–C(14) 118.4(6) C(8)–C(10)–N(3)–C(13) -2(1) C(12)–C(10)–N(3)–C(14) 5(1) O(3A)–C(4A)–C(5A)–O(4A) 3.2(9) O(3B)–C(4B)–C(5B)–O(4B) -1(1) C(6A)–C(5A)–O(4A)–C(7A) 6.3(9) C(6B)–C(5B)–O(4B)–C(7B) -7(1) Table 3 Hydrogen-bond parameters in co-crystals 6 and 8 donor–H acceptor H,A/A° D,A/A° D–H,A (°) symmetry operator co-crystal 6 O(3A)–H(3O) O(3B) 1.384(5) 2.471(7) 167(4) x y z N(2)–H(2N) O(3B) 1.751(4) 2.663(8) 137(3) x y z N(2)–H(2N) O(4B) 2.256(5) 3.175(8) 140(4) x y z co-crystal 8 N(3)–H(3) O(4) 1.86(4) 2.653(4) 162(3) x y z N(3)–H(3) O(3) 1.97(5) 2.818(4) 122(2) x y z O(5)–H(51) O(4) 2.53(4) 3.046(4) 170(3) x y z between the phenoxide anion and pyridinium cation the heterometric trimer is held together by three intermolecular hydrogen bonds as listed in Table 3.The phenoxide oxygen atom O(3B) of the anionic phenol forms an OMH,O hydrogen bond to the hydroxy group of the neutral phenol. The bond distances of O(3B),H(3O) [1.384(5) A° ] and O(3B),O(3A) [2.471(7) A° ] indicate that this is a very strong OMH,O hydrogen bond.22 The phenoxide oxygen atom also accepts a hydrogen H(2N) from the pyridinium cation to form a strong NMH,O hydrogen bond O(3B),N(2) [2.663(8) A° ] which is considerately shorter than most intraand inter-molecular N,O hydrogen-bond distances (2.85 A° ).21 The protonated hydrogen also forms a weak NMH,O hydrogen bond to the methoxy oxygen atom O(4B) of the anionic phenol with an O(4B),N(2) distance of 3.175(8) A° . Crystal structure analysis of 8 As with the structure of 6 8 is an organic salt composed of one molecule of 2-methoxy-4-nitrophenoxide (the anion) one molecule of 4-pyrrolidinylpyridinium (the cation) and one molecule of water.The water molecule is disordered and one Fig. 3 The crystal structure and atomic numbering of 8. Hydrogen of its hydrogen atoms was not located. The atomic numbering bonds are indicated by dashed lines. scheme of 8 is shown in Fig. 3. Selected bond distances bond angles and torsion angles of the co-crystal are listed in Table 4. The two aromatic planes of the phenoxide anion and the transfer the contribution of the quinoid form II shown in Fig. 2 increases in the structure of the phenoxide ion as pyridinium cation are slightly twisted from each other by a dihedral angle of 7.6°.The occurrence of proton transfer is reflected by the shorter bond lengths of C(4)MO(4) [1.280(4) A° ] and C(1)MN(1) [1.427(4) A° ] (Table 4) compared indicated by the short N(3)MH(3) bond length [0.82(4) A° ] which is shorter than the corresponding N(2)MH(2) bond to those of C(4A)MO(3A) [1.360(8) A° ] and C(1A)MN(1A) [1.47(1) A° ] of the neutral phenol in 6 (Table 2). Because of length [1.092(7) A° ] found in 6. As a consequence of the proton J. Mater. Chem. 1997 7(5) 713–720 717 Table 4 Selected bond distances (A° ) bond angles (degrees) and torsion angles (degrees) in 8 C(4)–O(4) 1.280(4) C(1)–N(1) 1.427(4) C(3)–O(3) 1.366(3) C(7)–O(3) 1.429(4) C(1)–C(2) 1.395(4) C(1)–C(6) 1.393(5) C(2)–C(3) 1.370(4) C(3)–C(4) 1.438(4) C(5)–C(6) 1.367(5) C(4)–C(5) 1.415(4) C(8)–N(3) 1.339(5) C(12)–N(3) 1.339(4) C(8)–C(9) 1.350(4) C(11)–C(12) 1.355(4) C(9)–C(10) 1.426(4) C(10)–C(11) 1.415(4) C(13)–N(2) 1.464(4) C(16)–N(2) 1.476(4) C(10)–N(2) 1.323(4) N(3)–H(3) 0.82(4) O(1)–N(1)–O(2) 121.6(3) C(3)–O(3)–C(7) 116.6(2) C(3)–C(4)–O(4) 121.3(2) C(5)–C(4)–O(4) 122.9(3) C(10)–N(2)–C(13) 124.5(2) C(10)–N(2)–C(16) 124.2(2) C(13)–N(2)–C(16) 111.0(3) O(3)–C(3)–C(4)–O(4) -0.9(4) C(2)–C(3)–O(3)–C(7) 3.7(4) C(9)–C(10)–N(2)–C(13) -2.8(4) C(11)–C(10)–N(2)–C(16) 2.9(4) the contribution of the quinoid form the bond lengths of whether the phenol co-crystals are organic salts by solid-state IR study only.C(2)MC(3) [1.370(4) A° ] and C(5)MC(6) [1.367(5) A° ] are noticeably shorter than the other CMC bond lengths in Another possible way to predict formation of organic salts is the pKa dierence (DpKa) between organic acids and bases.the quinoid ring C(1)MC(2) [1.395(4) A° ] C(3)MC(4) [1.438(4) A° ] C(1)MC(6) [1.393(5) A° ] and C(4)MC(5) When DpKa is large enough the acid–base pairs may form organic salts. Based on the study of various benzoic acid– [1.415(4) A° ]. Like the structure of 6 the phenoxide and pyridinium ions pyridine complexes Johnson and Rumon have found that a DpKa of 3.7 is large enough to allow the carboxylic proton to in 8 are held together by ionic attractions as well as hydrogenbonding interactions. The hydrogen H(3) of the pyridinium transfer to the pyridine nitrogen of the bases.23 The calculated pKa for the cyanophenols nitrophenols and the protonated ion forms a bifurcated NMH,O hydrogen bond to the phenoxide oxygen atom O(4) and the methoxy oxygen atom pyridines used in this work are given in Table 5.24 The pKa dierences (DpKa) between the protonated pyridine (base) and O(3) of the anionic phenol [O(4),N(3) 2.653(4) A° and O(3),N(3) 2.818(4) A° ].The phenoxide oxygen atom also phenols (acid) were calculated and are listed in Table 6 (column 2). Since both 6 and 8 are ionic co-crystals and since accepts a proton H(51) from the water molecule to form an OMH,O hydrogen bond [O(4),O(5) 3.046(4) A° ]. the pKa dierence for the two co-crystals is 2.95 phenol– Determination of ionic or neutral phenol co-crystals Table 6 The DpKa stoichiometric ratio and SHG activity of co- Once the phenol–pyridine pairs complexed the resulting co- crystals 1–22 crystals were investigated to see whether proton transfer occurred in the co-crystals.As discussed previously in the stoichiometric SHG co-crystal DpKaa ratiob organic salt eciencyc structure analyses of 6 and 8 both of them are ionic cocrystals. The melting points of these two ionic co-crystals 1 2.95 251 yesd 2×urea (99–101 °C for 6 and 68–77 °C for 8) are not as high as those 2 2.95 151 e 0 of typical inorganic salts. (The broad melting range of 8 is 3 -1.49 251f g 0 probably due to dehydration of the co-crystal.) Thus the 4 -1.07 151 g 0 melting point measurement may not be a suitable method for 5 -1.89 151 g 0 6 2.95 251 yesh 2×urea determining the formation of ionic phenol co-crystals. Solid- 7 2.95 251 yesd 0 state IR has been used previously to determine whether benzoic 8 2.95 15151 yesh 2×urea acid–pyridine complexes are organic salts mainly based on 9 1.83 151 g 1×urea the comparison of n(CNO) between the carboxylic acid group 10 1.83 251 e 0 and the carboxylate group.23 Since there are only two crystal 11 1.83 151 g 0 structures available in this study it would be dicult to predict 12 2.95 15151 yesd 0 13 2.95 15151 yesd 0 14 2.95 15151 yesd 0 15 2.14 151 g 0 Table 5 The calculated pKa for phenol and pyridinium compoundsa 16 2.14 151 g 0 17 2.14 151 g 0 compound pKa 18 2.14 151 g 0 19 6.07 151 yesd 0 4-nitrophenol 7.15 2-methoxy-4-nitrophenol 7.15 20 6.07 151 yesd 0 21 3.00 251 yesd 0 2-methoxy-5-nitrophenol 8.27 4-cyanophenol 7.96 22 3.00 151 e 0 4-hydroxy-3-methoxybenzonitrile 7.96 2,4-dinitrophenol 4.03 aDpKa=pKa (base)-pKa (acid) where base is protonated pyridine and acid is phenol.bThe stoichiometric ratio of co-crystals is expressed 4-dimethylamino pyridinium 10.10 by the order (acid-15base5acid-2).cThe powder SHG eciency is based on the urea standard. dProposed by the consideration of 4-pyrrolidinopyridinium 10.10 4-phenyl pyridinium 5.26 DpKa2.95 and containing a third molecule except 19 and 20 which have a DpKa=6.07. eThese may be organic salts based on either trans-1,2-bis(4-pyridyl)ethylenium 6.08 trans-1,2-bis(4-pyridyl)ethanium 5.67 DpKa2.95 or containing a third molecule. fThe 251 ratio may not indicate the formation of salts because the base has two identical 4-(methylamino)pyridinium 10.15 pyridine nitrogen atoms. gIt is dicult to determine whether the compounds are ionic or neutral co-crystals. hConfirmed by crystal aThe pKa for the phenols and protonated pyridines (pyridinium) were calculated from ref.24. structure determination. 718 J. Mater. Chem. 1997 7(5) 713–720 Table 7 Analysis of the distribution of non-centrosymmetric structures for 1–22 from Table 6 number of number of compound non-centrosymmetricity data set analysis compounds with positive SHG (%) I all co-crystals 22 4 18 II co-crystals contain a third moleculea 10 3 30 III co-crystals are organic saltsb 10 3 30 aReflected by their 251 or 151 stoichiometric ratio (column 3 of Table 6). bThe co-crystals include 1 6–8 12–14 and 19–21 (column 4 of Table 6). pyridine pairs that have DpKa greater than 2.95 are expected Conclusion to form ionic co-crystals. Cyanophenols and nitrophenols are good hydrogen-bond There is only one good hydrogen-bond donor (the hydroxy donors and thus they can complex with various hydrogen- group) and one good hydrogen-bond acceptor (the pyridine N bond acceptors (pyridines) to form co-crystals.The phenol– atom) available in the phenol–pyridine pairs. The co-crystals pyridine pairs can form neutral or ionic co-crystals depending are expected to have a 151 stoichiometric ratio of phenol and on the acidity of the phenols and the basicity of the pyridines. pyridine. When proton transfer occurs between the hydrogen- When the pKa dierence of the pyridine and phenols is equal bond donor and acceptor the phenol–pyridine pair would to or greater than 2.95 the hydroxy proton of the phenols form organic salts that are composed of a phenoxide anion may transfer to the pyridine nitrogen atom of the pyridines and a pyridinium cation. As observed in the structures of 6 and form phenoxide–pyridinium salts (or ionic co-crystals).and 8 the phenoxide anion tends to accept a third molecule Crystal structure analyses revealed that 6 and 8 are ionic co- by hydrogen bonds suggesting that the anion is a very strong crystals and both crystallize in non-centrosymmetric structures. hydrogen-bond acceptor. In this case the stoichiometric ratio As anticipated the phenoxide and pyridinium ions in the co- of the co-crystals is more than 151 (e.g. 251 for 6 and 15151 crystals are held together by both ionic and hydrogen-bonding for 8). Thus the stoichiometric ratio of the co-crystals could interactions. Upon proton transfer the phenoxide anion tends be used as a guideline to predict whether proton transfer has to accept a third molecule via hydrogen bonds; thus the occurred in the co-crystals.Stoichiometric ratios determined resulting ionic co-crystals often have greater than 151 stoichio- from 1H NMR study for these co-crystals are given in Table 6 metric ratio. Preliminary results suggest that this type of co- (column 3). Based on the considerations of DpKa and stoichio- crystal particularly the ionic co-crystals may have a higher metric ratio 1 6–8 12–14 and 21 are expected to be organic chance of forming non-centrosymmetric structures than the salts since their DpKa is equal or larger than 2.95 and since normal achiral organic compounds. they contain a third molecule (Table 6 column 4). Although the stoichiometric ratio of 19 and 20 is 151 these two cocrystals are expected to be organic salts because their DpKa We gratefully acknowledge financial support from the Oce (6.07) is much higher than 2.95.In addition 2 10 and 22 are of Naval Research and the Byrn/Zografi joint project for study proposed as organic salts because they have either DpKa of the eect of water on the molecular mobility of pharmagreater than 2.95 or contain a third molecule. For rest of the ceutical solids. We would also like to thank Dr Isaac co-crystals it would be dicult to predict whether they are Ghebre-Sellasie (Pharmaceutical Research Warner-Lambert organic salts or not. Company) for his assistance in preparing this manuscript. Analysis of the distribution of non-centrosymmetric structures References for 1–22 1 Materials for Nonlinear Optics ed. S. R. Marder J. E. Sohn and Since the second-order NLO property is sensitive to the G.D. Stucky ACS Symp. Ser. 1991 455. 2 D. J. Williams Angew. Chem. Int. Ed. Engl. 1984 23 690. orientation of molecules in the solid-state medium,2 SHG 3 R. Rytel G. F. Lipscomb M. Stiller J. Thackara and measurements are useful for characterizing whether the phenol A. T. Ticknor in Nonlinear Optical Eects in Organic Polymers ed. co-crystals have non-centrosymmetric structures. The results J. Messier F. Kajzar P. Prasad and D. Ulrich Kluwer Academic of the SHG measurements are given in Table 6 (the last Publishers Dordrecht 1988 pp. 277–289. column). Of the twenty-two co-crystals studied 1 6 8 and 9 4 J. Zyss and J. L. Oudar Phys. Rev. 1982 A26 2016. have non-centrosymmetric crystal structures because they exhi- 5 R. J. Twieg and K. Jain in Nonlinear Optical Properties of Organic and Polymeric Materials ed.D. J. Williams ACS Symp. Ser. 1983 bit positive SHG activity. The non-centrosymmetric structures 233 57. of 6 and 8 were confirmed by X-ray structure determination 6 I. C. Paul and D. Y. Curtin Acc. Chem. Res. 1973 7 223. Pna21 (orthorhombic) for 6 and Cc (monoclinic) for 8. For the 7 G. R. Meredith in Nonlinear Optical Properties of Organic and twenty-two co-crystals the frequency of forming non-centro- Polymeric Materials ed. D. J. Williams ACS Symp. Ser. 1983 symmetric structures is 18% (set I of Table 7) which is slightly 233 27. higher than the value of 11% quoted for achiral organic 8 S. R. Marder J. W. Perry and W. P. Schaefer Science 1989 245 626. compounds in general.25 When only the co-crystals containing 9 J. Zyss and G. Berthier J.Chem. Phys. 1982 77 3635. a third molecule were considered 30% of the them have non- 10 T. W. Panunto Z. Urbanczyk-Lipkowska R. Johnson and centrosymmetric structures (set II of Table 7). Interestingly M. C. Etter J. Am. Chem. Soc. 1987 109 7786. this frequency is identical to that for the expected ionic 11 M. C. Etter and G. M. Frankenbach Chem. Mater. 1989 1 10. co-crystals (set III of Table 7). Although the number of the 12 C. B. Aakeroy P. B. Hitchcock B. D. Moyle and K. R. Seddon co-crystals studied here may not be large enough to make J. Chem. Soc. Chem. Commun. 1989 1856. 13 S. Kurtz and T. T. Perry J. Appl. Phys. 1968 39 3798. generalisations this type of phenol–pyridine co-crystals 14 J. P. Wibaut and F. W. Broekman Recl. T rav. Chim. Pays-Bas appears to show a higher frequency of forming non-centrosym- 1961 80 309.metric structures particularly for the ionic co-crystals than 15 N. Walker and D. Stuart Acta Crystallogr. 1983 A39 158. the normal achiral organic compounds. Thus the approach of 16 C. J. Gilmore J. Appl. Crystallogr. 1984 17 42. combining both ionic and hydrogen-bonding interactions may 17 P. T. Beurskens DIRDIF an Automatic Procedure for Phase Extension and Refinement of Dierence Structure Factors be useful for designing NLO materials. J. Mater. Chem. 1997 7(5) 713–720 719 Technique Report 1984/1 Crystallography Laboratory 21 R. J. Taylor and O. Kennard Acc. Chem. Res. 1984 17 320. 22 A. Novak Struct. Bonding (Berlin) 1974 18 177. Toernooiveld The Netherlands 1984. 18 R. Steward E. R. Davison and W. T. Simpson International T able 23 S.L. Johnson and K. A. Rumon J. Phys. Chem. 1965 69 74. 24 D. D. Perrin B. Dempsey and E. P. Serjeant in pKa prediction for for X-ray Crystallography Kynoch Press Birmingham 1974 vol. IV pp. 202–207. Organic Acids and Bases Chapman and Hall New York 1981. 25 M. C. Etter and K. S. Huang Chem. Mater. 1992 4 824. 19 TEXSAN TEXRAY Structure Analysis Package Molecular Structure Corporation 3200A Research Forest Drive The Woodlands TX 77381 USA 1985. Paper 6/04311J; Received 20th June 1996 20 L. J. Bellamy in T he Infrared Spectra of Complex Molecules Chapman and Hall New York vol. 2 1980. 720 J. Mater. Chem. 1997 7(5) 713–720

ISSN:0959-9428    

DOI:10.1039/a604311J

出版商:RSC    

年代:1997

数据来源: RSC

摘要:

Structural investigation of radiation grafted and sulfonated poly(vinylidene fluoride), PVDF, membranes Sami Hietala,a Svante Holmberg,b Milja Karjalainen,c Jan Na�sman,b Mikael Paronen,a Ritva Serimaa,c Franciska Sundholm*a and Sakari Vahvaselka�c aL aboratory of Polymer Chemistry, University of Helsinki, PB 55, FIN-00014 Helsinki, Finland bDepartment of Polymer T echnology, A° bo Akademi University, Porthansgatan 3-5, FIN-20500 A° bo, Finland cDepartment of Physics, University of Helsinki, PB 9, FIN-00014 Helsinki, Finland Radiation grafted and sulfonated poly(vinylidene fluoride), PVDF, membranes have been studied by thermal analysis and X-ray diraction to determine the changes in membrane crystallinity and structure during preparation.Commercial PVDF films were irradiated with an electron beam, grafted with styrene and finally sulfonated. Both the crystallinity and the size of the crystallites of PVDF decrease in the grafting reaction.A further decrease in crystallinity is observed in the sulfonation reaction. The residual crystallinity of PVDF was considerable (10–20%) even in membranes which had been subjected to severe reaction conditions.These results can be explained by assuming that the grafting takes place mainly in the amorphous region of the PVDF, and close to the surfaces of the crystals, but that grafts do not penetrate into the crystals. The proton conductivity of the grafted and sulfonated PVDF membranes reached values comparable to those of Nafion membranes. Membranes with high protonic conductivity are potentially values measured for Nafion were reached.It was found that the proton conductivity of the new membranes depended on useful as separators and electrolytes in electrochemical cells several factors, among them the crystallinity of the matrix such as fuel cells. Among the first proton-conducting mempolymer, and of the resulting graft, respectively. Residual branes used in fuel cells were sulfonated, crosslinked polycrystallinity was found in membranes which had been subject styrenes.1 In recent years Nafion films have been extensively even to the most severe reaction conditions, hence the resulting studied as proton-conducting membranes2 for use in fuel cells.membranes are inhomogeneous. Inherent inhomogeneity also Nafion is, however, a very expensive material, and other results from the preparation in which the macromolecules of materials are being sought.the original film act as the backbone with long branches In the development of new membranes, functionalisation of formed by the grafts. Because of the incompatibility of the two a preformed polymer material oers many advantages, in polymers the grafts are probably present as microdomains.particular with respect to the price and the mechanical proper- The bulk properties of the modified film, in this case most ties of the matrix material. Functionalisation is possible by importantly the proton conductivity, will depend on polymer– grafting. Among many grafting techniques available, radiationpolymer and polymer–water interactions, and on the distri- initiated grafting is one of the most useful.3,4 The electron bution and the phase separation of the grafts, as well as on beam pre-irradiation method is useful due to its simplicity and the crystallinity.17 short irradiation times, which makes the method attractive for The introduction of the styrene grafts and of the hydrophilic scale-up production.Radiation grafting methods have been sulfonic acid groups into the strongly hydrophobic PVDF extensively studied.3–5 matrix produces a very complex structure with the sulfonated Many studies have been reported on grafting dierent grafts mainly in the amorphous region of the PVDF.14 In the monomers onto fluorine-containing polymers by the radiation- present investigation the melting behaviour and the crystal- induced grafting technique.6–14 In order to obtain proton linity of the matrix material, the styrene grafted membranes, exchange membranes, styrene is usually grafted onto a fluoro- and of the sulfonated membranes, respectively, have been polymer and the graft copolymer is subsequently functionalised studied with thermal analysis, and with X-ray diraction, in with sulfonic acid groups.15 Because of the chemical stability order to evaluate the role of the grafting conditions, and of needed, the backbone polymers have been limited to the fully the sulfonation, on the microstructure of the membrane.The fluorinated polymers poly(tetrafluoroethylene), PTFE, or poly- matrix polymer is a non-porous film of PVDF, and the (tetrafluoroethylene-co-hexafluoropropylene), FEP.Styrene resulting styrene grafted and sulfonated membranes are proton together with divinylbenzene, DVB, has been grafted onto conducting in wet conditions. FEP followed by sulfonation. In fuel cells these membranes show good performance and stability.6,7 Poly(vinylidene fluoride), PVDF, is an interesting alternative, which apart from good stability also has a competitive price.Partially fluorinated Experimental membranes that fulfil the stability requirements in electrolysis Preparation of the grafted and sulfonated membranes, PVDF- cells have been described.16 g-PSSA We have recently reported the preparation and characterisation of proton-conducting membranes with porous poly- PVDF was supplied by Goodfellow as 80 mm thick films made (vinylidene fluoride), PVDF, films as matrices.14 The prep- by melt processing. Proton-conducting films were prepared in aration involves the pre-irradiation of PVDF film with an a three-step procedure.14 Pre-irradiated films (electron beam electron beam, graft polymerisation of styrene monomer into under nitrogen gas, dose 100 kGy) were grafted in a solution the matrix and sulfonation of the grafted film.The degree of of styrene in tetrahydrofuran (80520, v/v) at 67°C, or in a grafting and degree of sulfonation are controlled by the reaction solution of styrene in toluene (50550, v/v) at 70°C, to various degrees. The membrane with degree of grafting (d.o.g.)=2% time and conditions. Proton conductivities corresponding to J. Mater. Chem., 1997, 7(5), 721–726 721was prepared at ambient temperature, reaction time 4 h.The measured with Mo-Ka1 radiation, monochromated with a siliconmonochromator crystal in the incident beam, at ambient grafting time varied from 0.5 to 12 h at 67°C. The grafted membranes were dried to constant mass and the d.o.g. was temperature and at 170°C with annealing in vacuum. A h–h diractometer with a high-temperature chamber was used in determined gravimetrically, using eqn.(1) symmetrical transmission mode. The heated sample was meas- d.o.g.=[(m1-m0 )/m0]×100% (1) ured again after 12 h at ambient temperature and after a few weeks. The scattered intensities were measured with a scintil- where m0 is the mass of the original film and m1 the mass of lation counter. The SAXS measurements were made with a the grafted membrane.The grafted membranes were fully sealed copper anode fine-focus X-ray tube, used in point sulfonated [degree of sulfonation (d.o.s.)=100%] with 0.5 M focusing mode. The Cu-Ka1 radiation was monochromated by chlorosulfonic acid in 1,2-dichlorethane for 24 h. The sulfonmeans of a nickel filter and a totally reflecting glass block ation occurred mainly at the para position of the phenyl rings.18 (Huber small angle chamber 701).The scattered radiation was measured in the horizontal direction (beam width) by a linear Thermal analysis by dierential scanning calorimetry (DSC) one-dimensional position-sensitive proportional counter Thermograms of films with d.o.g. between 2 and 100% were (MBraun OED-50M). The space between the sample and the measured with a Perkin-Elmer DSC-7 calorimeter with a detector was evacuated to 0.05 mmHg using 13 mm polyimide heating rate of 20°C min-1 in the temperature range -50 to foils as X-ray windows.The scattering distance was 150 mm. 200 °C. The samples were dried for 7 days at 50°C in vacuo A narrow vertical slit was used before the sample to minimise prior to the measurement, after drying the membranes were background scatterine beam height at the sample was kept in a desiccator at ambient temperature. The measurement approximately 10 mm. The primary beam was narrow was repeated three times for each sample in order to determine (FWHM<0.002 A° -1). Together with the detector height pro- the eect of annealing on the shape of the melting peak, and file, the FWHM of the instrumental function in the vertical on the melting temperature.The overall percentage crystallinity directionwas about 0.044 A° -1. The magnitude of the scattering of the sample (Xc) was evaluated from the heat of fusion of vector is defined as k=(4p/l)sinh, where h is half of the the sample [eqn. (2)] scattering angle, and l is the wavelength.The smallest achievable k was 0.015 A° -1. The experimental SAXS curves were Xc(%)=DHf/DH100×100% (2) corrected for absorption, air scattering and experimental where DHf is the measured enthalpy of fusion of the sample, smearing due to the vertical instrument function. DH100 is the enthalpy of fusion of the 100% crystalline sample=104.7 J g-1.19 Results and Discussion Because of the amorphous nature of the polystyrene the measured enthalpy of fusion can be attributed to the PVDF PVDF films have been chosen as matrix materials for func- crystallites in the membranes.The mass fraction of PVDF in tionalised membranes due to their high reactivity in electron the grafted membranes is given by eqn. (3) beam irradiation grafting. Our previous results showed that very ecient grafting and sulfonation can be achieved in WPVDF=mPVDF/(mPVDF+mstyr) (3) porous PVDF films, and considerable proton conductivity is where mPVDF is the mass of the PVDF, mstyr the mass of obtained.14 The grafting and sulfonation of non-porous films polystyrene grafts and WPVDF the mass fraction of PVDF.The of PVDF requires longer reaction times, and the d.o.g.in these crystallinity of the PVDF is then obtained [eqn. (4)] films is therefore generally lower. Several membranes of nonporous PVDF grafted with styrene were prepared. The degree XPVDF(%)=DHf/WPVDF×100% (4) of grafting varied between 2 and 100%. The grafted membranes The following relationship [eqn. (5)] was used for the grafted were fully sulfonated,14,18 and the proton conductivity was and sulfonated membranes measuredwith ac impedance spectroscopy.The proton conductivity at room temperature increased with increasing d.o.g. WPVDF= mPVDF mPVDF+mstyr+mSO-3 (5) reaching values around 100 mS cm-1 for d.o.g. around 100%. The chemical and the electrochemical characterisation of these where mSO-3 is the mass of sulfonic acid groups. The mass of membranes will be published elsewhere.19,20 In the following sulfonic acid groups was calculated from the amount of styrene.we report the eects of the reaction conditions in the membrane The crystallinity of PVDF in the PVDF-g-PSSA membranes preparation on the melting behaviour and crystallinity of the was then calculated according to eqn. (4). PVDF-g-PSSA membranes with non-porous matrices.The melting behaviour of pure PVDF has been extensively Thermogravimetry (TG) studied.21,22 The partial crystallinity of this material has been described as the result of two factors; the polymer is inherently The thermal degradation of the PVDF-g-PSSA membranes easy to crystallise because of the simple structure of the was measured with a Mettler TG-50 thermobalance in the repeating unit, but complete crystallisation is inhibited due to temperature range 50–250 °C with a heating rate of the presence of head-to-head bonds in the polymer backbone. 10°C min-1. Samples of 5–8 mg were dried in vacuo at 50°C The crystalline phase in a solution-cast PVDF film shows a for 18 h prior to measurement. double endotherm in the melting region characteristic of a form crystals.Results from DSC measurements at dierent X-Ray diraction measurements heating rates and electron microscopy indicated that the two endotherms were caused by a temperature dependent bimodal The PVDF-g-PSSA membranes were measured at ambient temperature in symmetrical transmission and reflection geo- distribution of lamellar thicknesses. The lower melting peak is attributed to melting of crystals originally present, and the metries with Cu-Ka radiation (1.542 A° ), using a h–2h diractometer. The Cu-Ka radiation was monochromated with higher melting peak is attributed to melting of crystals reorganised on heating.According to Nakagawa and Ishida19 lamellar a quartz monochromator in the incident beam. To reduce the eects of preferred orientation the sample was rotated during thickening occurs in the larger crystals during the heating process, and the smaller crystals melt out.The material formed the measurement. The background was measured separately and subtracted from the intensity curves of the membranes. by the partial melting is proposed to recrystallise in the cooling process. This is proposed to be the reason for the double The intensity curves were also corrected for absorption and background due to air scattering.The original PVDF film was endotherm observed in the melting region of the a form on 722 J. Mater. Chem., 1997, 7(5), 721–726Fig. 1 Melting thermograms, second scan, of pre-irradiated styrene Fig. 2 Cooling thermograms, second scan, of pre-irradiated styrene grafted PVDF membranes, heating rate 20 °C min-1.From bottom: grafted PVDF membranes, cooling rate 20°C min-1. From bottom: original PVDF, d.o.g. 23%, 54%, 83% and 100%, respectively. original PVDF, d.o.g. 23%, 54%, 83% and 100%, respectively. repeated heating.20 In this study X-ray scattering was measured Results are collected in Fig. 3 and 4. It is seen that the from samples of PVDF at ambient temperature and at 170°C.crystallinity of the PVDF matrix is only slightly aected by A change in the WAXS pattern was observed, and the new the grafting process, and the decrease in crystallinity at higher pattern was still observed from the sample after 24 h at room d.o.g. is a few per cent, see Fig. 4. Thus we conclude that the temperature. The original pattern of the monoclinic a form irradiation has not formed reactive radical sites inside the was recorded after the sample had been standing for several crystallites.The decrease in the overall crystallinity of the weeks.23 The conclusion is drawn that heat annealing produces grafted membranes is therefore mostly an eect of dilution, see a crystalline phase which is not thermally stable at room Fig. 3. It is also noteworthy that the degree of crystallisation, temperature. as calculated from the enthalpies of fusion, does not change The thermal behaviour of the PVDF film, the grafted with repeated heating cycles, hence the ordered part of PVDF membranes, and of the PVDF-g-PSSA membranes, respectively, were studied by dierential scanning calorimetry and thermogravimetry. The thermal behaviour of the styrene grafted membranes will be discussed first.Typical melting thermograms of the second scans from these samples are presented in Fig. 1. Melting thermograms were recorded in three subsequent scans. The glass transition of PVDF at -40 °C24 could not be detected in the grafted PVDF. A bimodal melting endotherm is seen in all the thermograms. The shape of the bimodal melting endotherm changed when the samples were reheated so that the area under the lower melting peak increased and the area under the higher melting peak decreased. The change in the total enthalpy of fusion in subsequent scans of PVDF, and of grafted membranes, was very small.The melting temperatures were the same in the dierent samples irrespective of d.o.g., and of number of heating cycles, and they were ca. 163 and 167 °C. Fig. 3 Overall crystallinity of membranes as a function of the d.o.g. The shape of the bimodal melting peak of the grafted Styrene grafted PVDF membranes from calorimetric measurements membranes resembles that of pure PVDF, and changes slightly (%), and from WAXS measurements (&), PVDF-g-PSSA membranes with d.o.g. The melting range broadens as the d.o.g.increases. from calorimetric measurements (#), and from WAXS measurements ($). Calorimetric data from second scan. The melting peak of the reorganised thicker lamellae decreases with increasing d.o.g. Hence, the lamellar growth in the melting region is inhibited with formation of the polystyrene grafts. The conclusion is drawn that the grafting takes place in the entire amorphous region, also in areas very close to the crystallite surfaces of the lamellae. Cooling thermograms of the grafted PVDF membranes are shown in Fig. 2. The crystallisation temperature decreases slightly with increasing d.o.g., possibly indicating mixing of polystyrene grafts into the crystallisable part of the PVDF melt. However, the polystyrene grafts are not miscible with the matrix PVDF polymer, most probably they form a separated phase within the amorphous region of the styrene grafted PVDF.7 A very weak transition around 100 °C24 which can be attributed to the glass transition of polystyrene, can be observed in the first DSC scans of grafted PVDF membranes with high d.o.g.The transition is obscured in calorimetric studies of the grafts with low d.o.g.Fig. 4 Crystallinity of the PVDF moiety in membranes as a function because of the annealing eect of the sample drying process of the d.o.g. Styrene grafted PVDF membranes from calorimetric (7 days at 70°C) immediately before the measurement. measurements (%), and from WAXS measurements (&), PVDF-g- The crystallinity of the grafted PVDF and the PVDF-g- PSSA membranes from calorimetric measurements (#), and from WAXS measurements ($).Calorimetric data from second scan. PSSA membranes was calculated from the enthalpies of fusion. J. Mater. Chem., 1997, 7(5), 721–726 723is changing very little in the grafting process. This is in The well known Debye formula was used, since the crystal size is small. By fitting the intensities of crystalline PVDF and accordance with results obtained by Gupta and Scherer25 in investigations of proton-conducting membranes prepared amorphous material to the experimental intensity curves, the crystallinity of the polymer is obtained as the ratio of the by c-irradiation induced styrene grafting of perfluorinated polymer films.integrals of the intensity of the amorphous component and the studied sample.Owing to the uncertainty in the determination The crystallinity of the PVDF, of the grafted PVDF membranes, and of the PVDF-g-PSSA membranes was studied of the model intensity of the amorphous material, however, the precision of the crystallinity is only 10%. By this method with wide angle X-ray scattering (WAXS) measurements. The estimation of the crystallinity from the X-ray diractionpattern a crystallinity of 44% was obtained for the PVDF film using integration limits of 102h50°. is based on the assumption that the measured scattering intensity is a linear combination of the intensities from the Another method to determine the crystallinity from X-ray data is to present the scattering intensity of the crystalline crystalline and from the amorphous regions.26 The X-ray diraction patterns are shown in Fig. 6 and 7, and the results material as a sum of Gaussians or other suitable functions.28 The sum of these functions and a chosen background is fitted of the crystallinity calculations are included in Fig. 3 and 4. The intensity curve from totally amorphous polystyrene has to the intensity curve. We applied this method by calculating the powder diraction pattern of PVDF by presenting the been used as the amorphous background, and the intensity of crystalline PVDF is calculated from the atomic coordinates.27 reflections as Gaussians, and adjusting their widths to match the experimental ones.29 For PVDF this method gave a crystallinity of 35%, and for the calculated intensity of perfectly crystalline PVDF a crystallinity of ca. 60% was obtained. This model underestimates the diuse background due to the small crystal size, and thus also the crystallinity. The results of the analysis of the crystallinity from WAXS measurements lead us to believe that the first described method is the more accurate, and this method is therefore used in the present investigation. The diraction patterns measured by reflection and transmission geometries do not match perfectly due to preferred orientation of the crystallites, Fig. 5.The 020 and 110 reflections are very intense in the diraction pattern measured with reflection geometry, but weak when measured with transmission geometry. On the other hand, the 021 and 002 reflections are intense in the pattern measured with transmission geometry but weak when measured with reflection geometry. The precision in crystallinity is ca. 10% for pure PVDF. The size of the PVDF crystallites in the membranes was determined from theWAXS measurements. In untreated PVDF film, the measured average crystal size is around 15 nm, which is larger than the size of the crystallites in the porous PVDF membranes, 11 nm.14 With grafting a decrease in crystal size is observed, at d.o.g.around 50% the crystallite size is 10 nm in grafted membranes, and at 100% d.o.g. it is ca. 8 nm; Fig. 5 WAXS intensity curves of PVDF membranes. The intensity sulfonation does not alter the crystallite size. curves from top: measured in symmetrical reflection geometry, in Thus, the WAXS results indicate decrease both in crystal- symmetrical transmission geometry, and calculated intensity curve for linity and in crystallite size with increasing d.o.g.The decrease the a form. Fig. 7 WAXS intensity curves of PVDF-g-PSSA membranes. From Fig. 6 WAXS intensity curves of styrene grafted PVDF membranes. From top: d.o.g. 23%, 54%, 83% and 100%, respectively. top: d.o.g. 23%, 54%, 83% and 100%, respectively. 724 J.Mater. Chem., 1997, 7(5), 721–726Fig. 8 Enthalpy of fusion of styrene grafted PVDF membranes as a function of the d.o.g. The line denotes values of the enthalpy of fusion, Fig. 9 Melting thermograms, second scan, of PVDF-g-PSSA memassuming that no decrease in crystallinity takes place during the branes. Heating rate 20 °C min-1. From bottom: original PVDF, d.o.g.grafting. 23%, 54%, 83% and 100%, respectively. in crystallinity is steeper as measured from X-ray diraction data than from calorimetric measurements, Fig. 3 and 4. The change in melting enthalpy corresponds to a dilution eect by the grafts in the membrane at low d.o.g., but with increasing d.o.g. the crystallinity seems to decrease more, Fig. 8. The straight line represents the enthalpy of fusion of the grafted membranes vs.the d.o.g. when the decrease in crystallinity is caused by dilution only. A similar eect has been reported for grafting of polyethylene, where crystal disruption is detected with d.o.g. of over 30%.30 Results reported by Gebel et al.31 are somewhat contradictory. They studied styrene grafted and sulfonated PVDF membranes prepared after c-irradiation.The irradiation dose was relatively low, 1–100 kGy, generally lower doses than used in the present investigation, and the matrix material consisted Fig. 10 Cooling thermograms, second scan, of PVDF-g-PSSA mem- of only 25 mm thick films. The authors concluded from results branes. Cooling rate 20 °C min-1. From bottom: original PVDF, d.o.g. of WAXS and SAXS measurements that the grafting reac- 23%, 54%, 83% and 100%, respectively.tion takes place in amorphous zones without disturbing the crystalline texture. An explanation of the discrepancy between results in the up on sulfonation. The crystallisation in the PVDF-g-PSSA membranes occurs at lower temperatures than in grafted present work from calorimetric measurements and from X-ray measurements is that X-ray scattering has its origin from samples, indicating the restrictive impact of the polyelectrolyte chains on the mobility and ability to crystallise of the unsubsti- highly ordered and crystalline areas in the PVDF membranes.In layers close to the crystallite surfaces the grafting has caused tuted PVDF blocks in the chain. A bimodal crystallisation endotherm is clearly seen at d.o.g.ca. 20%. We conclude that a decrease in the lamellar thickness, but chain ends and loops present are not ordered enough to contribute to the melting crystallisable parts of the PVDF chains are present in two dierent surroundings at this d.o.g. which can be detected enthalpy registered in the calorimetric measurements. Owing to uncertainties in the estimation of the melting enthalpies of calorimetrically. A portion of the PVDF crystallises close to the surface of the membranes where the graft density is large, the partly crystalline polymers, and the base line in the thermograms, the precision in crystallinity measurements by while the remainder crystallises in the inner parts of the membrane, into which the sulfonated grafts have not yet thermal analysis is relatively low, around 10%. From these observations, the conclusion is drawn that the grafting reaction penetrated. Using energy dispersive X-ray analysis we have shown that polystyrene sulfonic acid grafts become evenly occurs mainly in the amorphous regions of the PVDF, but can occur very close to the surfaces of the crystallites in the distributed throughout the membrane at d.o.g.>ca. 30%.14 An exothermic enthalpy change over a wide temperature matrix polymer so breaking up partly ordered chain structures. The styrene monomer has not acted as a solvent of PVDF, range (100–200 °C) is detected in the first heating scan of the PVDF-g-PSSA membranes and this eect is very pronounced but has diused into the mobile amorphous regions to the reactive sites.at high d.o.g. The behaviour was not found in the second and subsequent scans, but was found again after storage of the Next we turn to the eect of sulfonation on the thermal behaviour and the crystallinity of the membranes. Typical membranes under ambient conditions. This behaviour is probably due to the hygroscopic properties of the PVDF-g-PSSA second scan heating thermograms of PVDF-g-PSSA membranes are seen in Fig. 9, the corresponding cooling scans are membranes, which bind water to the sulfonic acid groups. Not even drying at 100 °C for long periods could remove all the seen in Fig. 10. The melting behaviour of the PVDF-g-PSSA membranes resembles that of the polystyrene grafted PVDF water. This was also confirmed with thermogravimetry, where the mass loss from the PVDF-g-PSSA membranes (dried at membranes. The bimodal melting peak is detectable in all the samples.The form of the endotherm diers slightly with the 50°C for 18 h) was constant between 100 and 200 °C. Loss of sulfonic acid groups or degradation of polystyrene is not intensity of the first melting peak increasing relative to the second.There is a considerable decrease in the total melting probable in this temperature range, as shown in thermogravimetry measurements. enthalpy as the d.o.g. increases. This means that possible order remaining after the grafting reaction has been partly broken WAXS measurements were made of the PVDF-g-PSSA J. Mater. Chem., 1997, 7(5), 721–726 725formation of slightly ordered structures close to the lamellar surface of the PVDF which do not contribute to the melting enthalpy, but which are seen in X-ray diraction.Tero Lehtinen and Go�ran Sundholm are thanked for the electrochemical measurements and inspiring discussions. Matti Elomaa and Britta Mattsson are thanked for assistance with DSC measurements. The authors are indebted to several external sponsors for financial aid. S.Hietala and F. Sundholm wish in particular to thank the Nordic Energy Research Programme (NEFP) for funding and encouragement in this Nordic Cooperation. S. Holmberg, M. Paronen and F. Sundholm are indebted to the Academy of Finland for materials research funding (MATRA). References 1 L. W. Niedrach, W. T. Grubb, in Fuel Cells, ed. W. Mitchell, Academic Press, New York, 1963, ch. 3. 2 Fuel Cell Handbook, ed. A. J. Appleby and R. L. Foulkes, Van Nostrand, New York, 1989, pp. 277–295. 3 L. Mandelkern, in T he Radiation Chemistry ofMacromolecules, ed. Fig. 11 SAXS intensity curves for (bottom to top) PVDF film, and M. Dole, Academic Press, New York, 1972, vol. 1, ch. 13. hydrated PVDF-g-PSSA membranes with d.o.g. 23%, 54% and 100%, 4 D.O. Geymer, in T he Radiation Chemistry of Macromolecules, ed. respectively M. Dole, Academic Press, New York, 1973, vol. 1, ch. 1. 5 J. L. Garnett, Radiat. Phys. Chem., 1979, 14, 79. 6 B. Gupta, F. N. Bu�chi and G. G. Scherer, J. Polym. Sci., Part A: Polym. Chem., 1994, 32, 1931. membranes at ambient temperature, see Fig. 7. The crystallinity 7 B. Gupta and G. G. Scherer, Chimia, 1994, 48, 127.determined from these measurements is included in Fig. 3 and 8 F. Vigo, G. Capanelli, C. Uliana and S.Munari, Desalination, 1981, 4. It is evident that the sulfonation further decreases the 36, 63. crystallinity. The decrease is more dramatic at low d.o.g. and 9 E.-S. A. Hegazy, I. Ishigaki, A. M. Dessouki, A. Rabie and levels out at higher d.o.g. This reduction in crystallinity corre- J.Okamoto, J. Appl. Polym. Sci., 1982, 27, 535. lates quite well with the increase in electric conductivity found 10 A. Niemo�ller and G. Ellinghorst, Makromol. Chem., 1987, 148, 1. 11 F. Bu�chi, B. Gupta, M. Rouilly, P. C. Hauser, A. Chapiro and for PVDF-g-PSSA membranes when d.o.g. is above 20–30%.21 G. G. Scherer, Proc. 27th IECEC Conf., San Diego, August 1992, We assume that the number of grafts attached to the PVDF Society of Automotive Engineers, San Diego, CA, 1992, vol. 3, is rather similar in all these membranes, since the radiation p. 419. dose is the same in all cases. As the polystyrene graft chains 12 N. Betz, A. Le Moe�l, J. P. Durand, E. Balanzat and C. Darnez, become longer the eect of the hydrophilic sulfonic acid groups Macromolecules, 1992, 25, 213. 13 A. Elmidaoui, A. T. Cherif, J. Brunea, F. Duclert, T. Cohen and on the crystallites of PVDF becomes smaller. The size of the C. Gavach, J. Membr. Sci., 1992, 67, 263. crystallites is not much aected by the sulfonation, as judged 14 S. Holmberg, T. Lehtinen, J. Na�sman, D. Ostrovskii, M. Paronen, from WAXS results. This seems to indicate that whole crystal- R.Serimaa, F. Sundholm, G. Sundholm, L. Torell and M. Torkkeli, lites are destroyed in the sulfonation. This may be a result of J. Mater. Chem., 1996, 6, 1309. the extremely strong interaction between the hydrophilic sul- 15 G. G. Scherer, Ber. Bunsen-Ges. Phys. Chem., 1990, 94, 1008. 16 G. G. Scherer, E. Killer and D. Grman, Int. J. Hydrog. Energy, fonic acid groups and the hydrophobic PVDF matrix, which 1992, 17, 115.can cause crystal disruption.30 The complex, multiphase com- 17 A. Chapiro and A. M. Jedrychowska-Bonamour, Eur. Polym. J., position of the hydrated PVDF-g-PSSA membranes, consisting 1984, 20, 1079. of crystallites and amorphous regions of PVDF, and hydro- 18 M. V. Rouilly, R. Ko�tz, O. Haas, G. G. Scherer and A. Chapiro, philic sulfonated regions of polystyrene is further evidenced J.Membr. Sci., 1993, 81, 89. with SAXS diraction measurements, Fig. 11. A distinct feature 19 T. Lehtinen, F. Sundholm, G. Sundholm, P. Bjo�rnbom and M. Bursell, Electrochim. Acta, submitted. in the SAXS pattern is the interference maximum at k 20 S. Hietala, M. Karjalainen, T. Lehtinen, D. Ostrovskii, # 0.24 A° -1, corresponding to a Bragg spacing of 2.5 nm, M.Paronen, R. Serimaa, F. Sundholm and G. Sundholm, Appl. which indicates supramolecular order. In our previous report14 Macromol. Chem. Phys., submitted. this was attributed to water–sulfonic acid clusters, supposed 21 K. Nagakawa and Y. Ishida, J. Polym. Sci., Phys. Edn., 1973, 11, to be the active sites for the proton conduction in the mem- 2153. 22 E. Benedetti, S. Catanorchi, A. D’Alessio, G. Moggi, P. Vergamini, branes. The detailed structure and the state of water in the M. Pracella and F. Ciardelli, Polym. Int., 1996, 41, 35. hydrated PVDF-g-PSSA is currently being studied in our 23 R. Serimaa and S. Vahvaselka�, unpublished work. laboratory. 24 J. Brandrup and E. H. Immergut, Polymer Handbook, Wiley, New York, 3rd edn., 1989, p. VI-226 and VI-258. 25 B. Gupta and G. G. Scherer,Makromol. Chem., 1993, 210, 151. Conclusion 26 F. J. Balta�-Calleja and C. G. Vonk, X-Ray Scattering of Synthetic Polymers, Elsevier, Amsterdam, 1989, pp. 175–204. PVDF-g-PSSA membranes prepared by the three-step pro- 27 R. Hasegawa, Y. Takahashi, Y. Chatani and H. Tade of electron beam irradiation, grafting and sulfonation J., 1972, 3, 600. undergo changes in crystallinity under the reaction conditions. 28 L. D. Majdanac, D. Poleti and M. J. Teodorovic, Acta Polymerica, The crystallinity of the membranes decreases not only by the 1991, 42, 351. dilution eect of the grafted polystyrene in the amorphous 29 W. Kraus and G. Nolze, the computer program Powder Cell, version 1.8. regions of the PVDF matrix, but also by the disordering eect 30 R. Y. M. Huang and P. J. F. Kanitz, J. Appl. Polym. Sci., 1969, of the grafts of the lamellar surfaces. Crystal disruption can 13, 669. not be excluded. The sulfonation of the grafted membranes 31 G. Gebel, E. Ottomani, J. J. Allegraud, N. Betz and A. Le Moe�l, further deteriorates the crystalline region of PVDF. Slightly Nucl. Instrum.Meth. Phys. Res. B, 1995, 105, 145. diering values of the crystallinity are obtained by calorimetric Paper 6/07675K; Received 12th November, 1996 analysis and by X-ray diraction analysis. This is due to 726 J. Mater. Chem., 1997, 7(5), 721–7

ISSN:0959-9428    

DOI:10.1039/a607675k

出版商:RSC    

年代:1997

数据来源: RSC

摘要:

Kinetic and mechanistic aspects of copper(II ) coordination to bis-N,N¾-(salicylidene)- 1,2-diaminoethane-based hydrogel polymer membranes, and the permeation of cations through them Andrew J. Hall and J. David Miller* T he Speciality Materials Research Group, T he Department of Chemical Engineering and Applied Chemistry, Aston University, Aston T riangle, Birmingham, UK B4 7ET A range of hydrophilic membranes composed of copolymers of bis-N,N¾-(5-vinylsalicylidene)-1,2-diaminoethane with 2-hydroxyethylmethacrylate have been synthesised.Over a period of approximately 2 h these membranes coordinate copper(II) ions from aqueous solution to yield tetradentate species in a first-order process. However, only a small fraction of the potentially ligating sites are fully used in this way.Kinetic studies of the interactions with the nitrate, chloride and sulfate salts of copper(II) are described and a detailed mechanism is proposed. Molecular rotations at the ligand site are suggested to be the rate determining steps of the overall process. Values of individual rate and equilibrium constants have been determined, and shown to be consistent with the equivalent data found for simpler ligands involved in reactions in homogeneous solutions. The permeation of the nitrates of CoII, NiII and CuII through membranes of these copolymers is also described.Due to the slow rates of complex formation, the ligand sites have no significant eect on either the permeability of the salts through the membrane, or the time lags before salt passage is detected. For comparison purposes, permeation data for the passage of the nitrate of the substitutionally inert [Cr(H2O)6 ]3+ ion through 4-methyl-4¾-vinyl-2,2¾-bipyridyl-containing membranes are also reported.There is widespread interest in syntheses and separations using Abruna et al.,23 while bis-N,N¾-(5-vinylsalicylidene)-1,2-diaminoethane (H2dvsalen) was prepared by the method of Wul functional polymers.1 Advantage is often taken of the interand Akelah.24 actions which occur between metal ions and these func- Satisfactory characterisation of both monomers was tionalised polymers.This can range from the use of polymer achieved. Elemental analysis for vbpy gave C, 79.3; H, 6.3; N, supported complex catalysts through separation and hydrome- 14.0% (expected values being 79.6, 6.2 and 14.3%). 1H NMR tallurgy,2 and polymer-modified electrodes,3,4 to the investi- (CDCl3 ) d: 8.49/8.55 (H 6/6¾), 8.20/8.35 (H 5/5¾), 7.08/7.27 (H gation of more purely physical processes.5 The closely related 3/3¾), 6.67/6.7/6.72/6.76 (NCH2: position 4¾), 5.45/5.56, studies of the transport of metal ions through hydrogel mem- 6.01/6.07 (NCH: position 4¾) and 2.39 (CH3: position 4). 13C branes has been reported by other workers.6,7 Here we touch NMR (CDCl3) d: 155.8/156.8 (C 2/2¾), 148.8/149.3 (C 6/6¾), on transport but mainly focus on some aspects of the coordi- 145.7/148.1 (C 4/4¾), 121.9/124.7 (C 5/5¾), 118.7 (NCH2: pos- nation process itself.ition 4¾), 118.4/120.5 (C 3/3¾) and 21.1 (CH3 : position 4). Studies involving two dierent hydrogel copolymers, The other functionalised-ligand monomer, H2dvsalen, was incorporating well known ligand groups, are described here.prepared in 85% yield from its immediate precursor 5-vinyl The first, 2,2¾-bipyridyl (bipy), has been extensively studied as salicylaldehyde (Found: C, 72.4; H, 5.2. Calc. for H2dvsalen: a monomeric ligand in solution.8–11 It has also been func- C, 73.0; H, 5.4%). 1H NMR (CDCl3 ) d: 13.25 (br, H 9/9¾), 8.33 tionalised, usually as 4-methyl-4¾-vinyl-2,2¾-bipyridine (vbpy) (H 7/7¾), 7.35/7.36/7.38/7.39 (H 3/3¾), 7.22/7.23/7.24 (H 5/5¾), and incorporated into copolymers. Members of this research 6.88/6.91 (H 6/6¾), 6.56/6.59/6.61/6.65 (NCH: positions 4/4¾), group have concentrated on hydrogel copolymers, and used 5.09/5.12, 5.54/5.59 (NCH2 : positions 4/4¾) and 3.93 (H 8/8¾).them in coordination studies.12–17 The other ligand group 13C NMR (CDCl3 ) d: 166.4 (C 7/7¾), 160.8 (C 1/1¾), 135.6 described in this paper is bis-N,N¾-(salicylidene)-1,2-diamino- (NCH: positions 4/4¾), 130.0 (C 5/5¾), 129.4 (C 4/4¾), 118.3 (C ethane (H2salen). Again it has been studied as a discrete 2/2¾), 111.6 (NCH2: positions 4/4¾) and 59.7 (C 8/8¾).The ligand18–20 and as a functionalised component of copolymers, melting point of our product is 167–170 °C, which compares e.g. using the 5,5¾-divinyl derivative (H2dvsalen) which we have with the literature value of 168–170 °C. examined here.21,22 The main emphasis in this report is upon kinetic and Membrane fabrication mechanistic studies involving the coordination of the Cu2+ ion, but we begin by briefly describing some relevant trans- Membranes of approximately 0.4mm unhydrated, but accu- port studies.rately measured, thickness were prepared by a previously described method.13 The desired mixture of monomers, also containing 0.5% AIBN and 1% EGDM, was purged with Experimental nitrogen, injected into a glass mould through a G22 syringe needle, and heated at 60°C for three days, followed by 2 h Monomers postcure at 90°C.The membrane was then extracted from the Optical grade 2-hydroxyethyl methacrylate (HEMA) was used mould, inspected and measured. The variations in thickness as supplied by Kelvin Lenses Limited. Ethylene glycol dimetha- measured at dierent positions on a sheet were found to be crylate (EGDM) cross-linking agent was obtained from BDH within 2% of the average value, while films showing defects and used without further purification. The free radical initiator such as pinholes were discarded at this stage.Membranes were azobisisobutyronitrile (AIBN) was obtained from Aldrich then left to hydrate in deionised water for 2 weeks, with daily and recrystallised before use. 4-Methyl-4¾-vinyl-2,2¾-bipyridine changes of water. The membranes were stored in deionised water until required. (vbpy) was prepared in an overall 30% yield by the method of J. Mater. Chem., 1997, 7(5), 727–732 727Kinetic and mechanistic aspects of copper(II ) coordination to bis-N,N¾-(salicylidene)- 1,2-diaminoethane-based hydrogel polymer membranes, and the permeation of cations through them Andrew J.Hall and J. David Miller* T he Speciality Materials Research Group, T he Department of Chemical Engineering and Applied Chemistry, Aston University, Aston T riangle, Birmingham, UK B4 7ET A range of hydrophilic membranes composed of copolymers of bis-N,N¾-(5-vinylsalicylidene)-1,2-diaminoethane with 2-hydroxyethylmethacrylate have been synthesised.Over a period of approximately 2 h these membranes coordinate copper(II) ions from aqueous solution to yield tetradentate species in a first-order process. However, only a small fraction of the potentially ligating sites are fully used in this way. Kinetic studies of the interactions with the nitrate, chloride and sulfate salts of copper(II) are described and a detailed mechanism is proposed.Molecular rotations at the ligand site are suggested to be the rate determining steps of the overall process. Values of individual rate and equilibrium constants have been determined, and shown to be consistent with the equivalent data found for simpler ligands involved in reactions in homogeneous solutions. The permeation of the nitrates of CoII, NiII and CuII through membranes of these copolymers is also described.Due to the slow rates of complex formation, the ligand sites have no significant eect on either the permeability of the salts through the membrane, or the time lags before salt passage is detected. For comparison purposes, permeation data for the passage of the nitrate of the substitutionally inert [Cr(H2O)6 ]3+ ion through 4-methyl-4¾-vinyl-2,2¾-bipyridyl-containing membranes are also reported.There is widespread interest in syntheses and separations using Abruna et al.,23 while bis-N,N¾-(5-vinylsalicylidene)-1,2-diaminoethane (H2dvsalen) was prepared by the method of Wul functional polymers.1 Advantage is often taken of the interand Akelah.24 actions which occur between metal ions and these func- Satisfactory characterisation of both monomers was tionalised polymers.This can range from the use of polymer achieved. Elemental analysis for vbpy gave C, 79.3; H, 6.3; N, supported complex catalysts through separation and hydrome- 14.0% (expected values being 79.6, 6.2 and 14.3%). 1H NMR tallurgy,2 and polymer-modified electrodes,3,4 to the investi- (CDCl3 ) d: 8.49/8.55 (H 6/6¾), 8.20/8.35 (H 5/5¾), 7.08/7.27 (H gation of more purely physical processes.5 The closely related 3/3¾), 6.67/6.7/6.72/6.76 (NCH2: position 4¾), 5.45/5.56, studies of the transport of metal ions through hydrogel mem- 6.01/6.07 (NCH: position 4¾) and 2.39 (CH3: position 4). 13C branes has been reported by other workers.6,7 Here we touch NMR (CDCl3) d: 155.8/156.8 (C 2/2¾), 148.8/149.3 (C 6/6¾), on transport but mainly focus on some aspects of the coordi- 145.7/148.1 (C 4/4¾), 121.9/124.7 (C 5/5¾), 118.7 (NCH2: pos- nation process itself.ition 4¾), 118.4/120.5 (C 3/3¾) and 21.1 (CH3 : position 4). Studies involving two dierent hydrogel copolymers, The other functionalised-ligand monomer, H2dvsalen, was incorporating well known ligand groups, are described here.prepared in 85% yield from its immediate precursor 5-vinyl The first, 2,2¾-bipyridyl (bipy), has been extensively studied as salicylaldehyde (Found: C, 72.4; H, 5.2. Calc. for H2dvsalen: a monomeric ligand in solution.8–11 It has also been func- C, 73.0; H, 5.4%). 1H NMR (CDCl3 ) d: 13.25 (br, H 9/9¾), 8.33 tionalised, usually as 4-methyl-4¾-vinyl-2,2¾-bipyridine (vbpy) (H 7/7¾), 7.35/7.36/7.38/7.39 (H 3/3¾), 7.22/7.23/7.24 (H 5/5¾), and incorporated into copolymers.Members of this research 6.88/6.91 (H 6/6¾), 6.56/6.59/6.61/6.65 (NCH: positions 4/4¾), group have concentrated on hydrogel copolymers, and used 5.09/5.12, 5.54/5.59 (NCH2 : positions 4/4¾) and 3.93 (H 8/8¾). them in coordination studies.12–17 The other ligand group 13C NMR (CDCl3 ) d: 166.4 (C 7/7¾), 160.8 (C 1/1¾), 135.6 described in this paper is bis-N,N¾-(salicylidene)-1,2-diamino- (NCH: positions 4/4¾), 130.0 (C 5/5¾), 129.4 (C 4/4¾), 118.3 (C ethane (H2salen).Again it has been studied as a discrete 2/2¾), 111.6 (NCH2: positions 4/4¾) and 59.7 (C 8/8¾). The ligand18–20 and as a functionalised component of copolymers, melting point of our product is 167–170 °C, which compares e.g.using the 5,5¾-divinyl derivative (H2dvsalen) which we have with the literature value of 168–170 °C. examined here.21,22 The main emphasis in this report is upon kinetic and Membrane fabrication mechanistic studies involving the coordination of the Cu2+ ion, but we begin by briefly describing some relevant trans- Membranes of approximately 0.4mm unhydrated, but accu- port studies.rately measured, thickness were prepared by a previously described method.13 The desired mixture of monomers, also containing 0.5% AIBN and 1% EGDM, was purged with Experimental nitrogen, injected into a glass mould through a G22 syringe needle, and heated at 60°C for three days, followed by 2 h Monomers postcure at 90°C. The membrane was then extracted from the Optical grade 2-hydroxyethyl methacrylate (HEMA) was used mould, inspected and measured. The variations in thickness as supplied by Kelvin Lenses Limited.Ethylene glycol dimetha- measured at dierent positions on a sheet were found to be crylate (EGDM) cross-linking agent was obtained from BDH within 2% of the average value, while films showing defects and used without further purification.The free radical initiator such as pinholes were discarded at this stage. Membranes were azobisisobutyronitrile (AIBN) was obtained from Aldrich then left to hydrate in deionised water for 2 weeks, with daily and recrystallised before use. 4-Methyl-4¾-vinyl-2,2¾-bipyridine changes of water. The membranes were stored in deionised water until required. (vbpy) was prepared in an overall 30% yield by the method of J.Mater. Chem., 1997, 7(5), 727–732 727Kinetic and mechanistic aspects of copper(II ) coordination to bis-N,N¾-(salicylidene)- 1,2-diaminoethane-based hydrogel polymer membranes, and the permeation of cations through them Andrew J. Hall and J. David Miller* T he Speciality Materials Research Group, T he Department of Chemical Engineering and Applied Chemistry, Aston University, Aston T riangle, Birmingham, UK B4 7ET A range of hydrophilic membranes composed of copolymers of bis-N,N¾-(5-vinylsalicylidene)-1,2-diaminoethane with 2-hydroxyethylmethacrylate have been synthesised.Over a period of approximately 2 h these membranes coordinate copper(II) ions from aqueous solution to yield tetradentate species in a first-order process.However, only a small fraction of the potentially ligating sites are fully used in this way. Kinetic studies of the interactions with the nitrate, chloride and sulfate salts of copper(II) are described and a detailed mechanism is proposed. Molecular rotations at the ligand site are suggested to be the rate determining steps of the overall process. Values of individual rate and equilibrium constants have been determined, and shown to be consistent with the equivalent data found for simpler ligands involved in reactions in homogeneous solutions.The permeation of the nitrates of CoII, NiII and CuII through membranes of these copolymers is also described. Due to the slow rates of complex formation, the ligand sites have no significant eect on either the permeability of the salts through the membrane, or the time lags before salt passage is detected.For comparison purposes, permeation data for the passage of the nitrate of the substitutionally inert [Cr(H2O)6 ]3+ ion through 4-methyl-4¾-vinyl-2,2¾-bipyridyl-containing membranes are also reported. There is widespread interest in syntheses and separations using Abruna et al.,23 while bis-N,N¾-(5-vinylsalicylidene)-1,2-diaminoethane (H2dvsalen) was prepared by the method of Wul functional polymers.1 Advantage is often taken of the interand Akelah.24 actions which occur between metal ions and these func- Satisfactory characterisation of both monomers was tionalised polymers.This can range from the use of polymer achieved. Elemental analysis for vbpy gave C, 79.3; H, 6.3; N, supported complex catalysts through separation and hydrome- 14.0% (expected values being 79.6, 6.2 and 14.3%). 1H NMR tallurgy,2 and polymer-modified electrodes,3,4 to the investi- (CDCl3 ) d: 8.49/8.55 (H 6/6¾), 8.20/8.35 (H 5/5¾), 7.08/7.27 (H gation of more purely physical processes.5 The closely related 3/3¾), 6.67/6.7/6.72/6.76 (NCH2: position 4¾), 5.45/5.56, studies of the transport of metal ions through hydrogel mem- 6.01/6.07 (NCH: position 4¾) and 2.39 (CH3: position 4). 13C branes has been reported by other workers.6,7 Here we touch NMR (CDCl3) d: 155.8/156.8 (C 2/2¾), 148.8/149.3 (C 6/6¾), on transport but mainly focus on some aspects of the coordi- 145.7/148.1 (C 4/4¾), 121.9/124.7 (C 5/5¾), 118.7 (NCH2: pos- nation process itself.ition 4¾), 118.4/120.5 (C 3/3¾) and 21.1 (CH3 : position 4). Studies involving two dierent hydrogel copolymers, The other functionalised-ligand monomer, H2dvsalen, was incorporating well known ligand groups, are described here. prepared in 85% yield from its immediate precursor 5-vinyl The first, 2,2¾-bipyridyl (bipy), has been extensively studied as salicylaldehyde (Found: C, 72.4; H, 5.2.Calc. for H2dvsalen: a monomeric ligand in solution.8–11 It has also been func- C, 73.0; H, 5.4%). 1H NMR (CDCl3 ) d: 13.25 (br, H 9/9¾), 8.33 tionalised, usually as 4-methyl-4¾-vinyl-2,2¾-bipyridine (vbpy) (H 7/7¾), 7.35/7.36/7.38/7.39 (H 3/3¾), 7.22/7.23/7.24 (H 5/5¾), and incorporated into copolymers. Members of this research 6.88/6.91 (H 6/6¾), 6.56/6.59/6.61/6.65 (NCH: positions 4/4¾), group have concentrated on hydrogel copolymers, and used 5.09/5.12, 5.54/5.59 (NCH2 : positions 4/4¾) and 3.93 (H 8/8¾).them in coordination studies.12–17 The other ligand group 13C NMR (CDCl3 ) d: 166.4 (C 7/7¾), 160.8 (C 1/1¾), 135.6 described in this paper is bis-N,N¾-(salicylidene)-1,2-diamino- (NCH: positions 4/4¾), 130.0 (C 5/5¾), 129.4 (C 4/4¾), 118.3 (C ethane (H2salen).Again it has been studied as a discrete 2/2¾), 111.6 (NCH2: positions 4/4¾) and 59.7 (C 8/8¾). The ligand18–20 and as a functionalised component of copolymers, melting point of our product is 167–170 °C, which compares e.g. using the 5,5¾-divinyl derivative (H2dvsalen) which we have with the literature value of 168–170 °C. examined here.21,22 The main emphasis in this report is upon kinetic and Membrane fabrication mechanistic studies involving the coordination of the Cu2+ ion, but we begin by briefly describing some relevant trans- Membranes of approximately 0.4mm unhydrated, but accu- port studies.rately measured, thickness were prepared by a previously described method.13 The desired mixture of monomers, also containing 0.5% AIBN and 1% EGDM, was purged with Experimental nitrogen, injected into a glass mould through a G22 syringe needle, and heated at 60°C for three days, followed by 2 h Monomers postcure at 90°C.The membrane was then extracted from the Optical grade 2-hydroxyethyl methacrylate (HEMA) was used mould, inspected and measured. The variations in thickness as supplied by Kelvin Lenses Limited.Ethylene glycol dimetha- measured at dierent positions on a sheet were found to be crylate (EGDM) cross-linking agent was obtained from BDH within 2% of the average value, while films showing defects and used without further purification. The free radical initiator such as pinholes were discarded at this stage. Membranes were azobisisobutyronitrile (AIBN) was obtained from Aldrich then left to hydrate in deionised water for 2 weeks, with daily and recrystallised before use. 4-Methyl-4¾-vinyl-2,2¾-bipyridine changes of water. The membranes were stored in deionised water until required. (vbpy) was prepared in an overall 30% yield by the method of J. Mater. Chem., 1997, 7(5), 727–732 727Kinetic and mechanistic aspects of copper(II ) coordination to bis-N,N¾-(salicylidene)- 1,2-diaminoethane-based hydrogel polymer membranes, and the permeation of cations through them Andrew J.Hall and J. David Miller* T he Speciality Materials Research Group, T he Department of Chemical Engineering and Applied Chemistry, Aston University, Aston T riangle, Birmingham, UK B4 7ET A range of hydrophilic membranes composed of copolymers of bis-N,N¾-(5-vinylsalicylidene)-1,2-diaminoethane with 2-hydroxyethylmethacrylate have been synthesised.Over a period of approximately 2 h these membranes coordinate copper(II) ions from aqueous solution to yield tetradentate species in a first-order process. However, only a small fraction of the potentially ligating sites are fully used in this way. Kinetic studies of the interactions with the nitrate, chloride and sulfate salts of copper(II) are described and a detailed mechanism is proposed.Molecular rotations at the ligand site are suggested to be the rate determining steps of the overall process. Values of individual rate and equilibrium constants have been determined, and shown to be consistent with the equivalent data found for simpler ligands involved in reactions in homogeneous solutions.The permeation of the nitrates of CoII, NiII and CuII through membranes of these copolymers is also described. Due to the slow rates of complex formation, the ligand sites have no significant eect on either the permeability of the salts through the membrane, or the time lags before salt passage is detected. For comparison purposes, permeation data for the passage of the nitrate of the substitutionally inert [Cr(H2O)6 ]3+ ion through 4-methyl-4¾-vinyl-2,2¾-bipyridyl-containing membranes are also reported.There is widespread interest in syntheses and separations using Abruna et al.,23 while bis-N,N¾-(5-vinylsalicylidene)-1,2-diaminoethane (H2dvsalen) was prepared by the method of Wul functional polymers.1 Advantage is often taken of the interand Akelah.24 actions which occur between metal ions and these func- Satisfactory characterisation of both monomers was tionalised polymers.This can range from the use of polymer achieved. Elemental analysis for vbpy gave C, 79.3; H, 6.3; N, supported complex catalysts through separation and hydrome- 14.0% (expected values being 79.6, 6.2 and 14.3%). 1H NMR tallurgy,2 and polymer-modified electrodes,3,4 to the investi- (CDCl3 ) d: 8.49/8.55 (H 6/6¾), 8.20/8.35 (H 5/5¾), 7.08/7.27 (H gation of more purely physical processes.5 The closely related 3/3¾), 6.67/6.7/6.72/6.76 (NCH2: position 4¾), 5.45/5.56, studies of the transport of metal ions through hydrogel mem- 6.01/6.07 (NCH: position 4¾) and 2.39 (CH3: position 4). 13C branes has been reported by other workers.6,7 Here we touch NMR (CDCl3) d: 155.8/156.8 (C 2/2¾), 148.8/149.3 (C 6/6¾), on transport but mainly focus on some aspects of the coordi- 145.7/148.1 (C 4/4¾), 121.9/124.7 (C 5/5¾), 118.7 (NCH2: pos- nation process itself.ition 4¾), 118.4/120.5 (C 3/3¾) and 21.1 (CH3 : position 4). Studies involving two dierent hydrogel copolymers, The other functionalised-ligand monomer, H2dvsalen, was incorporating well known ligand groups, are described here.prepared in 85% yield from its immediate precursor 5-vinyl The first, 2,2¾-bipyridyl (bipy), has been extensively studied as salicylaldehyde (Found: C, 72.4; H, 5.2. Calc. for H2dvsalen: a monomeric ligand in solution.8–11 It has also been func- C, 73.0; H, 5.4%). 1H NMR (CDCl3 ) d: 13.25 (br, H 9/9¾), 8.33 tionalised, usually as 4-methyl-4¾-vinyl-2,2¾-bipyridine (vbpy) (H 7/7¾), 7.35/7.36/7.38/7.39 (H 3/3¾), 7.22/7.23/7.24 (H 5/5¾), and incorporated into copolymers.Members of this research 6.88/6.91 (H 6/6¾), 6.56/6.59/6.61/6.65 (NCH: positions 4/4¾), group have concentrated on hydrogel copolymers, and used 5.09/5.12, 5.54/5.59 (NCH2 : positions 4/4¾) and 3.93 (H 8/8¾).them in coordination studies.12–17 The other ligand group 13C NMR (CDCl3 ) d: 166.4 (C 7/7¾), 160.8 (C 1/1¾), 135.6 described in this paper is bis-N,N¾-(salicylidene)-1,2-diamino- (NCH: positions 4/4¾), 130.0 (C 5/5¾), 129.4 (C 4/4¾), 118.3 (C ethane (H2salen). Again it has been studied as a discrete 2/2¾), 111.6 (NCH2: positions 4/4¾) and 59.7 (C 8/8¾). The ligand18–20 and as a functionalised component of copolymers, melting point of our product is 167–170 °C, which compares e.g.using the 5,5¾-divinyl derivative (H2dvsalen) which we have with the literature value of 168–170 °C. examined here.21,22 The main emphasis in this report is upon kinetic and Membrane fabrication mechanistic studies involving the coordination of the Cu2+ ion, but we begin by briefly describing some relevant trans- Membranes of approximately 0.4mm unhydrated, but accu- port studies. rately measured, thickness were prepared by a previously described method.13 The desired mixture of monomers, also containing 0.5% AIBN and 1% EGDM, was purged with Experimental nitrogen, injected into a glass mould through a G22 syringe needle, and heated at 60°C for three days, followed by 2 h Monomers postcure at 90°C.The membrane was then extracted from the Optical grade 2-hydroxyethyl methacrylate (HEMA) was used mould, inspected and measured. The variations in thickness as supplied by Kelvin Lenses Limited. Ethylene glycol dimetha- measured at dierent positions on a sheet were found to be crylate (EGDM) cross-linking agent was obtained from BDH within 2% of the average value, while films showing defects and used without further purification.The free radical initiator such as pinholes were discarded at this stage. Membranes were azobisisobutyronitrile (AIBN) was obtained from Aldrich then left to hydrate in deionised water for 2 weeks, with daily and recrystallised before use. 4-Methyl-4¾-vinyl-2,2¾-bipyridine changes of water.The membranes were stored in deionised water until required. (vbpy) was prepared in an overall 30% yield by the method of J. Mater. Chem., 1997, 7(5), 727–732 727Kinetic and mechanistic aspects of copper(II ) coordination to bis-N,N¾-(salicylidene)- 1,2-diaminoethane-based hydrogel polymer membranes, and the permeation of cations through them Andrew J. Hall and J.David Miller* T he Speciality Materials Research Group, T he Department of Chemical Engineering and Applied Chemistry, Aston University, Aston T riangle, Birmingham, UK B4 7ET A range of hydrophilic membranes composed of copolymers of bis-N,N¾-(5-vinylsalicylidene)-1,2-diaminoethane with 2-hydroxyethylmethacrylate have been synthesised. Over a period of approximately 2 h these membranes coordinate copper(II) ions from aqueous solution to yield tetradentate species in a first-order process.However, only a small fraction of the potentially ligating sites are fully used in this way. Kinetic studies of the interactions with the nitrate, chloride and sulfate salts of copper(II) are described and a detailed mechanism is proposed. Molecular rotations at the ligand site are suggested to be the rate determining steps of the overall process.Values of individual rate and equilibrium constants have been determined, and shown to be consistent with the equivalent data found for simpler ligands involved in reactions in homogeneous solutions. The permeation of the nitrates of CoII, NiII and CuII through membranes of these copolymers is also described.Due to the slow rates of complex formation, the ligand sites have no significant eect on either the permeability of the salts through the membrane, or the time lags before salt passage is detected. For comparison purposes, permeation data for the passage of the nitrate of the substitutionally inert [Cr(H2O)6 ]3+ ion through 4-methyl-4¾-vinyl-2,2¾-bipyridyl-containing membranes are also reported.There is widespread interest in syntheses and separations using Abruna et al.,23 while bis-N,N¾-(5-vinylsalicylidene)-1,2-diaminoethane (H2dvsalen) was prepared by the method of Wul functional polymers.1 Advantage is often taken of the interand Akelah.24 actions which occur between metal ions and these func- Satisfactory characterisation of both monomers was tionalised polymers.This can range from the use of polymer achieved. Elemental analysis for vbpy gave C, 79.3; H, 6.3; N, supported complex catalysts through separation and hydrome- 14.0% (expected values being 79.6, 6.2 and 14.3%). 1H NMR tallurgy,2 and polymer-modified electrodes,3,4 to the investi- (CDCl3 ) d: 8.49/8.55 (H 6/6¾), 8.20/8.35 (H 5/5¾), 7.08/7.27 (H gation of more purely physical processes.5 The closely related 3/3¾), 6.67/6.7/6.72/6.76 (NCH2: position 4¾), 5.45/5.56, studies of the transport of metal ions through hydrogel mem- 6.01/6.07 (NCH: position 4¾) and 2.39 (CH3: position 4). 13C branes has been reported by other workers.6,7 Here we touch NMR (CDCl3) d: 155.8/156.8 (C 2/2¾), 148.8/149.3 (C 6/6¾), on transport but mainly focus on some aspects of the coordi- 145.7/148.1 (C 4/4¾), 121.9/124.7 (C 5/5¾), 118.7 (NCH2: pos- nation process itself.ition 4¾), 118.4/120.5 (C 3/3¾) and 21.1 (CH3 : position 4). Studies involving two dierent hydrogel copolymers, The other functionalised-ligand monomer, H2dvsalen, was incorporating well known ligand groups, are described here. prepared in 85% yield from its immediate precursor 5-vinyl The first, 2,2¾-bipyridyl (bipy), has been extensively studied as salicylaldehyde (Found: C, 72.4; H, 5.2.Calc. for H2dvsalen: a monomeric ligand in solution.8–11 It has also been func- C, 73.0; H, 5.4%). 1H NMR (CDCl3 ) d: 13.25 (br, H 9/9¾), 8.33 tionalised, usually as 4-methyl-4¾-vinyl-2,2¾-bipyridine (vbpy) (H 7/7¾), 7.35/7.36/7.38/7.39 (H 3/3¾), 7.22/7.23/7.24 (H 5/5¾), and incorporated into copolymers.Members of this research 6.88/6.91 (H 6/6¾), 6.56/6.59/6.61/6.65 (NCH: positions 4/4¾), group have concentrated on hydrogel copolymers, and used 5.09/5.12, 5.54/5.59 (NCH2 : positions 4/4¾) and 3.93 (H 8/8¾). them in coordination studies.12–17 The other ligand group 13C NMR (CDCl3 ) d: 166.4 (C 7/7¾), 160.8 (C 1/1¾), 135.6 described in this paper is bis-N,N¾-(salicylidene)-1,2-diamino- (NCH: positions 4/4¾), 130.0 (C 5/5¾), 129.4 (C 4/4¾), 118.3 (C ethane (H2salen).Again it has been studied as a discrete 2/2¾), 111.6 (NCH2: positions 4/4¾) and 59.7 (C 8/8¾). The ligand18–20 and as a functionalised component of copolymers, melting point of our product is 167–170 °C, which compares e.g. using the 5,5¾-divinyl derivative (H2dvsalen) which we have with the literature value of 168–170 °C.examined here.21,22 The main emphasis in this report is upon kinetic and Membrane fabrication mechanistic studies involving the coordination of the Cu2+ ion, but we begin by briefly describing some relevant trans- Membranes of approximately 0.4mm unhydrated, but accu- port studies. rately measured, thickness were prepared by a previously described method.13 The desired mixture of monomers, also containing 0.5% AIBN and 1% EGDM, was purged with Experimental nitrogen, injected into a glass mould through a G22 syringe needle, and heated at 60°C for three days, followed by 2 h Monomers postcure at 90°C.The membrane was then extracted from the Optical grade 2-hydroxyethyl methacrylate (HEMA) was used mould, inspected and measured.The variations in thickness as supplied by Kelvin Lenses Limited. Ethylene glycol dimetha- measured at dierent positions on a sheet were found to be crylate (EGDM) cross-linking agent was obtained from BDH within 2% of the average value, while films showing defects and used without further purification. The free radical initiator such as pinholes were discarded at this stage.Membranes were azobisisobutyronitrile (AIBN) was obtained from Aldrich then left to hydrate in deionised water for 2 weeks, with daily and recrystallised before use. 4-Methyl-4¾-vinyl-2,2¾-bipyridine changes of water. The membranes were stored in deionised water until required. (vbpy) was prepared in an overall 30% yield by the method of J.Mater. Chem., 1997, 7(5), 727–732 727Kinetic and mechanistic aspects of copper(II ) coordination to bis-N,N¾-(salicylidene)- 1,2-diaminoethane-based hydrogel polymer membranes, and the permeation of cations through them Andrew J. Hall and J. David Miller* T he Speciality Materials Research Group, T he Department of Chemical Engineering and Applied Chemistry, Aston University, Aston T riangle, Birmingham, UK B4 7ET A range of hydrophilic membranes composed of copolymers of bis-N,N¾-(5-vinylsalicylidene)-1,2-diaminoethane with 2-hydroxyethylmethacrylate have been synthesised.Over a period of approximately 2 h these membranes coordinate copper(II) ions from aqueous solution to yield tetradentate species in a first-order process. However, only a small fraction of the potentially ligating sites are fully used in this way.Kinetic studies of the interactions with the nitrate, chloride and sulfate salts of copper(II) are described and a detailed mechanism is proposed. Molecular rotations at the ligand site are suggested to be the rate determining steps of the overall process. Values of individual rate and equilibrium constants have been determined, and shown to be consistent with the equivalent data found for simpler ligands involved in reactions in homogeneous solutions.The permeation of the nitrates of CoII, NiII and CuII through membranes of these copolymers is also described. Due to the slow rates of complex formation, the ligand sites have no significant eect on either the permeability of the salts through the membrane, or the time lags before salt passage is detected.For comparison purposes, permeation data for the passage of the nitrate of the substitutionally inert [Cr(H2O)6 ]3+ ion through 4-methyl-4¾-vinyl-2,2¾-bipyridyl-containing membranes are also reported. There is widespread interest in syntheses and separations using Abruna et al.,23 while bis-N,N¾-(5-vinylsalicylidene)-1,2-diaminoethane (H2dvsalen) was prepared by the method of Wul functional polymers.1 Advantage is often taken of the interand Akelah.24 actions which occur between metal ions and these func- Satisfactory characterisation of both monomers was tionalised polymers. This can range from the use of polymer achieved.Elemental analysis for vbpy gave C, 79.3; H, 6.3; N, supported complex catalysts through separation and hydrome- 14.0% (expected values being 79.6, 6.2 and 14.3%). 1H NMR tallurgy,2 and polymer-modified electrodes,3,4 to the investi- (CDCl3 ) d: 8.49/8.55 (H 6/6¾), 8.20/8.35 (H 5/5¾), 7.08/7.27 (H gation of more purely physical processes.5 The closely related 3/3¾), 6.67/6.7/6.72/6.76 (NCH2: position 4¾), 5.45/5.56, studies of the transport of metal ions through hydrogel mem- 6.01/6.07 (NCH: position 4¾) and 2.39 (CH3: position 4). 13C branes has been reported by other workers.6,7 Here we touch NMR (CDCl3) d: 155.8/156.8 (C 2/2¾), 148.8/149.3 (C 6/6¾), on transport but mainly focus on some aspects of the coordi- 145.7/148.1 (C 4/4¾), 121.9/124.7 (C 5/5¾), 118.7 (NCH2: pos- nation process itself. ition 4¾), 118.4/120.5 (C 3/3¾) and 21.1 (CH3 : position 4).Studies involving two dierent hydrogel copolymers, The other functionalised-ligand monomer, H2dvsalen, was incorporating well known ligand groups, are described here. prepared in 85% yield from its immediate precursor 5-vinyl The first, 2,2¾-bipyridyl (bipy), has been extensively studied as salicylaldehyde (Found: C, 72.4; H, 5.2. Calc. for H2dvsalen: a monomeric ligand in solution.8–11 It has also been func- C, 73.0; H, 5.4%). 1H NMR (CDCl3 ) d: 13.25 (br, H 9/9¾), 8.33 tionalised, usually as 4-methyl-4¾-vinyl-2,2¾-bipyridine (vbpy) (H 7/7¾), 7.35/7.36/7.38/7.39 (H 3/3¾), 7.22/7.23/7.24 (H 5/5¾), and incorporated into copolymers. Members of this research 6.88/6.91 (H 6/6¾), 6.56/6.59/6.61/6.65 (NCH: positions 4/4¾), group have concentrated on hydrogel copolymers, and used 5.09/5.12, 5.54/5.59 (NCH2 : positions 4/4¾) and 3.93 (H 8/8¾).them in coordination studies.12–17 The other ligand group 13C NMR (CDCl3 ) d: 166.4 (C 7/7¾), 160.8 (C 1/1¾), 135.6 described in this paper is bis-N,N¾-(salicylidene)-1,2-diamino- (NCH: positions 4/4¾), 130.0 (C 5/5¾), 129.4 (C 4/4¾), 118.3 (C ethane (H2salen). Again it has been studied as a discrete 2/2¾), 111.6 (NCH2: positions 4/4¾) and 59.7 (C 8/8¾). The ligand18–20 and as a functionalised component of copolymers, melting point of our product is 167–170 °C, which compares e.g. using the 5,5¾-divinyl derivative (H2dvsalen) which we have with the literature value of 168–170 °C. examined here.21,22 The main emphasis in this report is upon kinetic and Membrane fabrication mechanistic studies involving the coordination of the Cu2+ ion, but we begin by briefly describing some relevant trans- Membranes of approximately 0.4mm unhydrated, but accu- port studies. rately measured, thickness were prepared by a previously described method.13 The desired mixture of monomers, also containing 0.5% AIBN and 1% EGDM, was purged with Experimental nitrogen, injected into a glass mould through a G22 syringe needle, and heated at 60°C for three days, followed by 2 h Monomers postcure at 90°C. The membrane was then extracted from the Optical grade 2-hydroxyethyl methacrylate (HEMA) was used mould, inspected and measured. The variations in thickness as supplied by Kelvin Lenses Limited. Ethylene glycol dimetha- measured at dierent positions on a sheet were found to be crylate (EGDM) cross-linking agent was obtained from BDH within 2% of the average value, while films showing defects and used without further purification. The free radical initiator such as pinholes were discarded at this stage. Membranes were azobisisobutyronitrile (AIBN) was obtained from Aldrich then left to hydrate in deionised water for 2 weeks, with daily and recrystallised before use. 4-Methyl-4¾-vinyl-2,2¾-bipyridine changes of water. The membranes were stored in deionised water until required. (vbpy) was prepared in an overall 30% yield by the method of J. Mater. Chem., 1997, 7(5), 727–732 727

ISSN:0959-9428    

DOI:10.1039/a606999a

出版商:RSC    

年代:1997

数据来源: RSC

摘要:

Monitoring the growth of titanium oxide thin films by the liquid-phase deposition method with a quartz crystal microbalance Shigehito Deki, Yoshifumi Aoi, Yasuhiro Asaoka, Akihiko Kajinami and Minoru Mizuhata Department of Chemical Science & Engineering, Faculty of Engineering, Kobe University, Rokkodai-cho, Nada-ku, Kobe 657, Japan The quartz crystal microbalance (QCM) technique has been applied to investigate the formation of titanium oxide thin films by the liquid-phase deposition (LPD) method.A linear relationship was observed between the thickness measured by the QCM technique and that measured by direct observation with a scanning electron microscope, indicating that it is possible to monitor the growth of thin films from aqueous solution systems by the LPD method with the QCM technique.The concentration eects of free F-, H3BO3 and (NH4)2TiF6 on the film deposition rate are discussed. The liquid-phase deposition (LPD) method is a wet process equivalence is expressed as follows:4,11 for the formation of metal oxide thin films on substrates.1,2 By Df=-f02Dm/(rqNA) (4) using this method, it is possible to form metal oxide or where f0 is the fundamental frequency of the crystal, N is the hydroxide thin films on various kinds of substrates which are frequency constant (1.670×105 cm Hz for an AT-cut quartz) immersed only in the treatment solution.Metal oxide or and rq is the density of quartz (2.648 g cm-3). hydroxide thin films form through the ligand-exchange In the present paper, we report the growth of titanium oxide (hydrolysis) equilibrium reaction of metal–fluoro complex ionic thin films by the LPD method, which was monitored with the species and the F- consumption reaction of boric acid or QCM technique.We discuss the concentration eects of free aluminium metal as F- scavengers. In the treatment solution, F-, H3BO3 and (NH4 )2TiF6 in the treatment solution on thin following ligand-exchange (hydrolysis),an equilibrium reaction film formation by the LPD method.of the metal–fluoro complex ion is presumed: MFxn-+nH2O=MOn+xHF (1) Experimental The equilibrium reaction (1) is shifted to the right-hand side Treatment solutions for the LPD method by the addition of boric acid or aluminium metal as F- As parent solutions for the LPD method, (NH4)2TiF6 (Kishida scavengers, which readily react with F- and form more stable Chemical Co.Ltd.) and H3BO3 (Nacalai Tesque Inc.) were complex ions as follows: dissolved separately in distilled water at concentrations of 0.5 mol dm-3. These solutions were mixed in various composi- H3BO3+4HF=BF4-+H3O++2H2O (2) tions and used as the solution for deposition. For the purpose Al+6HF=H3AlF6+3/2H2 (3) of investigating the concentration eect of free F- in the treatment solution, an NH4F aqueous solution (Nacalai The addition of the F- scavenger leads to the consumption of Tesque Inc.) was added to the treatment solution at various F- ions and accelerates the ligand-exchange reaction (1).This concentrations. is a very simple process and does not require any special equipment, such as a vacuum system, a high frequency or QCM measurements voltage source.This process can be applied readily to the The quartz crystal used in the present study was biplanar, preparation of thin films on substrates which have complex circular AT-cut with a diameter of 13 mm. The crystal had a surface morphologies and large surface areas, and also multi- fundamental frequency of 3.58 MHz, and the approximate component oxide thin films can be formed by the addition of sensitivity for the frequency change was 35 ng cm-2 Hz-1 the selected metal ion to the treatment solution, because the [from eqn.(4)]. The surface of the crystal in contact with the LPD method is performed in an aqueous solution which is a solution was coated with a thin film of metallic gold (ca.typical homogeneous mixing system. So far, we have developed 100 nm) by vacuum evaporation. The crystal was mounted on and reported the preparation of titanium oxide thin films from a holder using a silicon sealant so as to expose one side of the a mixed solution of (NH4)2TiF6 and H3BO3 by the LPD crystal to the solution. method.3 A schematic diagram of the cell and quartz crystal apparatus The quartz crystal microbalance (QCM) is an extremely is shown in Fig. 1. The cell was made of acrylic resin. The sensitive sensor capable of measuring small mass changes in quartz crystal apparatus was immersed in the treatment solu- the nanogram range, because its oscillating frequency is tion for deposition, and the resonant frequency was measured changed even by a small amount of deposition on the crystal by a frequency counter (Advantest Inc., TR5823H) at intervals surface.4 Recently, it has been used widely as a mass-sensitive of 15 min at 30°C.detector in aqueous solution systems as well as in vacuum and in the gas phase. QCM techniques in solution systems have Results and Discussion been applied to the study of electrodeposition,5–7 ion fluxes in polymer films,8,9 intercalation phenomena10 and other interfa- Monitoring the growth of thin films with the QCM cial phenomena.The change in oscillating frequency of the quartz crystal (Df ) is proportional to the change in mass (Dm) Fig. 2 shows the relationship between the thickness of the film deposited by the LPD method as determined by SEM per unit area (A) of the material deposited on the crystal.The J. Mater. Chem., 1997, 7(5), 733–736 733Deposition of titanium oxide thin films by the LPD method The deposition reaction in the LPD method consists of the ligand-exchange (hydrolysis) equilibrium reaction of the metal– fluoro complex ion and the F- consumption reaction by the addition of boric acid as F- scavenger. For the ligand exchange (hydrolysis) of [TiF6]2- ion in aqueous solution, the following equilibrium scheme has been proposed.12,13 [TiF6 ]2-+nH2O=[TiF6-n(OH)n]2-n+nHF (7) The value of n is small, and the degree of hydrolysis decreases upon increasing the concentration of [TiF6]2-.12,13 On the other hand, H3BO3 readily reacts with F- as follows:14–16 H3BO3+HF=HBF(OH)3 (8) HBF(OH)3+HF=HBF2(OH)2+H2O (9) HBF2(OH)2+HF=HBF3OH+H2O (10) HBF3OH+HF=HBF4+H2O (11) Reaction steps (8)–(10) are rapid, and step (11) is slow.14–16 The addition of H3BO3 to the (NH4)2TiF6 solution accelerates the ligand-exchange (hydrolysis) reaction (7), i.e.reaction (7) is shifted to the right, owing to the reaction of H3BO3 with F-. Consequently, titanium oxide thin films form on the substrates upon dehydration of the [Ti(OH)6]2- species generated by the hydrolysis reaction of [TiF6]2-.Fig. 3(a) shows the relationship between the deposited film Fig. 1 Schematic diagram of the cell and quartz crystal apparatus. (a) Frequency counter, (b) dc power source, (c) quartz crystal oscillator, mass (m) and the reaction time. Fig. 3(b) shows the variation (d) oscillating circuit, (e) treatment solution and (f ) o-ring. The cell of the deposition rate (r) with reaction time.The deposition was kept at 30°C. rate (r) was determined from the gradient of Fig. 3(a). In this case, the deposition rate increased with reaction time up to 6 h, was constant to 9 h, and then decreased gradually. As shown in Fig. 3(b), there are three steps for the deposition rate: (i) increasing with time, (ii) a constant rate region and (iii) decreasing with time.From SEM observation of the Fig. 2 Relationship between film thickness determined from QCM(tQCM) and SEM (tSEM ) observation (tSEM ) and that measured by the QCM technique (tQCM). The value of tQCM was determined using the following equation. tQCM=Dm 1 Ad (5) where Dm is obtained from eqn. (4), A is the area of the QCM and d is the density of TiO2 (anatase; 3.90 g cm-3).As shown in Fig. 2, a good linear relationship was observed between tSEM and tQCM . It can be said, therefore, that the QCM technique is suitable for monitoring the growth of thin films by the LPD method. The value of tQCM was expressed as follows: tQCM=0.85 tSEM (6) The value of tQCM was lower than that of tSEM owing to the Fig. 3 Dependence of deposited film mass (a) and deposition rate (b) value of d used in eqn. (5) being larger than the real value, on reaction time. Concentration of (NH4)2TiF6 100 mmol dm-3, and of H3BO3 150 mmol dm-3. because the crystallinity of the deposited film was low. 734 J. Mater. Chem., 1997, 7(5), 733–736deposited film, several nucleation sites were observed on the rates were almost constant.The variation of induction period with the concentration of NH4F is shown in Fig. 4(b). The substrate at step (i), and the number of nucleation sites increased with the reaction time. At steps (ii) and (iii), growth induction period increased linearly with increasing concentration of NH4F. Based on these data, we concluded that free of the nuclei was observed.The decrease in the deposition rate at step (iii) is presumably due to the decrease in the reactant F- ions react with H3BO3 in the initial stage of the deposition reaction, i.e. the induction period. After that, F- ions content. Here we introduce some parameters. The reaction time at which the deposition rate was constant is tc. The coordinated to [TiF6-n(OH)n]2- react with H3BO3, and then titanium oxide deposition occurs.The film formation rate does deposited film mass at tc is mc. The deposition rate at tc is defined as the film formation rate, rc. The induction period for not depend on the concentration of free F-. deposition (i ) is defined as follows: Eects of the concentration of H3BO3 in the treatment solution i=tc-mc rc (12) Fig. 5(a) shows the relationship between the deposited film mass and the reaction time.The concentration of (NH4)2TiF6 These parameters are indicated on Fig. 3. was constant at 100 mmol dm-3 and that of H3BO3 was varied from 125.0 to 200.0 mmol dm-3. The deposited film mass Eects of the concentration of free F- in the treatment solution increased with increasing concentration of H3BO3. The film formation rate and the induction period are shown in Fig. 5(b) In the solution, [TiF6]2- ions partially hydrolyse and release as a function of the concentration of H3BO3. As can be seen, F-, although the value of n of eqn. (7) is small. Thus free F- the film formation rate increased and the induction period ions may exist in the solution at the initial stage. It is considered decreased with increasing H3BO3 concentration.Variation of that the H3BO3, as F- scavenger, preferentially reacts with the film formation rate was contrasted with that of the such free F- at the outset of the deposition reaction, therefore induction period. the induction period may be strongly correlated with the initial As mentioned, free F- ions exist in the [TiF6]2- aqueous content of free F-.In order to investigate the eects of the solution at equilibrium concentration. These free F- ions react concentration of free F-, anNH4F solution was added to the with H3BO3 during the induction period, and then F- ions treatment solution. As NH4F dissociates to NH4+ and F-, coordinated to the [TiF6-n(OH)n]2- complex ion react with the concentration of free F- in the treatment solution could H3BO3.The concentrations of free F- are constant for every be controlled by the addition of NH4F. Fig. 4(a) shows the solution in Fig. 5, and since the concentrations of (NH4)2TiF6 relationship between the reaction time and the deposited film are constant, the induction period decreased with increasing mass with various concentrations of NH4F in the treatment H3BO3 concentration. The decrease in the induction period solution.The deposited film mass decreased and the induction slowed when the concentration of H3BO3 was more than period increased with increasing concentration of NH4F, i.e. 162.5 mmol dm-3. As mentioned previously, reaction of F- increasing free F- concentration, whereas the film formation Fig. 5 (a) Relationship between deposited film mass and reaction time.Fig. 4 (a) Relationship between deposited film mass and reaction time. Concentration of NH4F: # 0, ' 5.0, % 10.0 and 1 15.0 mmol dm-3. Concentration of (NH4)2TiF6 50.0 mmol dm-3. Concentration of H3BO3 :# 125.0, ' 150.0, % 162.5, 1 175.0 and ×200.0 mmol dm-3. (b) Variation of induction period with concentration of NH4F. Concentration of (NH4 )2TiF6 100 mmol dm-3, and of H3BO3 (b) Variations of deposition rate (#) and induction period (') with concentration of H3BO3. 150 mmol dm-3. J. Mater. Chem., 1997, 7(5), 733–736 735concentration. The change of the induction period was smaller than that when the concentration of (NH4)2TiF6 was constant [Fig. 5(b)]. As mentioned above, the degree of hydrolysis of [TiF6 ]2- decreases with increasing (NH4)2TiF6 concentration.12 The change of concentration of free F- with the concentration of (NH4)2TiF6 may be small.Thus, the dependences of the induction period and tc on the concentration of (NH4)2TiF6 were small. On the other hand, the concentration of partially hydrolysed species of [TiF6]2- decreased with increasing (NH4)2TiF6 concentration, causing the decrease of the film formation rate with increasing (NH4)2TiF6 concentration.Conclusions The QCM technique has been applied to monitoring the growth of titanium oxide thin films by the LPD method. The concentration eects of free F-, (NH4)2TiF6 and H3BO3 in the treatment solution on film formation were studied. In the induction period, the initial stage of the deposition reaction, free F- ions in the treatment solution preferentially react with H3BO3.The induction period increased linearly upon increasing the concentration of added NH4F, indicating that the induction period depends on the concentration of free F-. The film formation rate, however, is independent of the concentration of free F-. The film formation rate increased and the induction period decreased with increasing H3BO3 concentration.When the concentration of H3BO3 was greater than 162.5 mol dm-3, the increase in the film formation rate and the decrease in the induction period slowed, suggesting that F- was consumed by H3BO3 in the slow step when the Fig. 6 (a) Relationship between deposited film mass and reaction concentration of H3BO3 was low; on the other hand, it was time.Concentration of H3BO3 200 mmol dm-3. Concentration of consumed in the fast steps when the concentration of H3BO3 (NH4)2TiF6: # 37.5, ' 50.0, % 75.0, 1 100.0, ×112.5 and was high. The film formation rate decreased with increasing +137.5 mmol dm-3. (b) Relationship between deposition rate (#) (NH4)2TiF6 concentration, and the dependence of the induc- and induction period (') with concentration of (NH4 )2TiF6 .tion period on the concentration of (NH4)2TiF6 was small. The QCM technique may be useful in monitoring the ion with H3BO3 can be divided into fast steps [eqn. (8)–(10)] formation of thin films by the LPD method for characterizing and a slow step [eqn. (11)]. It could be considered that when the kinetics of the growth of the films. the concentration of H3BO3 is higher than 162.5 mmol dm-3, free F- ions were consumed in the fast steps, therefore the induction period decreased gradually. On the other hand, the References film formation rate increased markedly with the H3BO3 concentration to 162.5 mmol dm-3, and then increased gradually. 1 H. Nagayama, H. Honda and H. Kawahara, J. Electrochem. Soc., The film formation rate is controlled by the reaction of F- 1988, 135, 2013. 2 A. Hishinuma, T. Goda, M. Kitaoka, S. Hayashi and ions coordinated to [TiF6-n(OH)n]2-. We consider that H. Kawahara, Appl. Surf. Sci., 1991, 48/49, 405. F- coordinated to [TiF6-n(OH)n]2- is consumed by H3BO3 3 S. Deki, Y. Aoi, O. Hiroi and A. Kajinami, Chem. L ett., 1996, 433. during the slow step [eqn. (11)] when the concentration of 4 G.Sauerbrey, Z. Phys., 1959, 155, 206. H3BO3 is lower than 162.5 mmol dm-3. On the other hand, 5 M. Seo, K. Yoshida, H. Takahashi and I. Sawamura, F- is consumed in the fast steps [eqn. (8)–(10)] when the J. Electrochem. Soc., 1992, 139, 3108. concentration of H3BO3 is higher than 162.5 mmol dm-3. 6 R. Beck, U. Pittermann and K. G.Weil, J. Electrochem. Soc., 1992, 139, 453.Therefore the dependence of the film formation rate on the 7 M. R. Deakin and O. R. Melroy, J. Electrochem. Soc., 1989, 136, H3BO3 concentration varied with the concentration range 349. of H3BO3. 8 J. H. Kaufman, K. K. Kanazawa and G. B. Street, Phys. Rev. L ett., 1984, 53, 2461. Eects of the concentration of (NH4 )2TiF6 in the treatment 9 D. Orata and D. A. Buttry, J. Am. Chem.Soc., 1987, 109, 3574. solution 10 H. K. Park and W. H. Smyrl, J. Electrochem. Soc., 1994, 141, L25. 11 K. K. Kanazawa and J. G. Gordon II, Anal. Chim. Acta, 1985, Fig. 6(a) shows the relationship between the deposited film 175, 99. mass and the reaction time for various concentrations of 12 R. H. Schmitt, E. L. Grove and R. D. Brown, J. Am. Chem. Soc., (NH4)2TiF6 . The concentration of H3BO3 was constant at 1960, 82, 5292. 13 Y. A. Buslaev, D. S. Dyer and R. O. Ragsdale, Inorg. Chem., 1967, 200 mmol dm-3. The deposited film mass decreased with 6, 2208. increasing concentration of (NH4)2TiF6. The tc values were 14 C. A. Wamser, J. Am. Chem. Soc., 1948, 70, 1209. almost constant at 3 h. Fig. 6(b) shows the film formation rate 15 R. E. Mesmer, K. M. Palen and C. F. Baes Jr., Inorg. Chem., 1973, and induction period as a function of the concentration of 12, 89. H3BO3. The film formation rate decreased monotonically with 16 C. A. Wamser, J. Am. Chem. Soc., 1951, 73, 409. increasing (NH4)2TiF6 concentration. On the other hand, the induction period increased slightly with increasing (NH4)2TiF6 Paper 6/07466I; Received 4th November, 1996 736 J. Mater. Chem., 1997, 7(5), 733–736

ISSN:0959-9428    

DOI:10.1039/a607466i

出版商:RSC    

年代:1997

数据来源: RSC

摘要:

Fabrication, characterization and photovoltaic study of a GaTSPc–CdS/TiO2 particulate film Jinghuai Fang,*a Jingwen Wu,a Lianyong Su,a Xiangyang Zhang,a Haifang Mao,b Yaochen Shena and Zuhong Lua aNational L aboratory of Molecular and Biomolecular Electronics, Southeast University, Nanjing, 210096 P.R. China bInstitute of Photographic Chemistry, Academic Sinica, Beijing 100101, P.R. China Quantum-sized cadmium sulfide (CdS) particles have been deposited onto the surface of a microporous titanium dioxide (TiO2) electrode by the chemical growth method, followed by modification with gallium tetrasulfophthalocyanine (GaTSPc) molecules to prepare a co-modified TiO2 electrode.Low bandgap semiconductor CdS particles and GaTSPc have dierent spectral ranges in the visible range. The co-modified TiO2 electrode eectively absorbed visible light.Experimental results showed that the sensitization of the TiO2 electrode with CdS and GaTSPc extended the optical absorption spectrum and photocurrent action spectrum into a wider visible range; moreover, the CdS/TiO2 coupled structure facilitated the photon-to-photocurrent conversion eciency. Recently there has been considerable interest in developing furnace in which the temperature was increased gradually to 450°C and then held at 450 °C for 30 min.After heat treatment, nanocrystalline microporous particulate films modified with organic dyes owing to the work of Gra�tzel and co-workers.1–3 the TiO2 film was strongly attached to the glass support. Using asimilar procedure to that described in ref. 1 and 8, a roughness The nanocrystalline semiconductor films possess highly porous structures and very large surface areas.Less than even a factor of about 80 was obtained for the TiO2 electrode. monolayer coating could completely absorb incident light in the spectral range of the dye absorbance; the energy conversion Fabrication of Q-CdS/TiO2 particulate film eciency could be improved. Subsequently, there have been In order to deposit CdS particles on the microporous TiO2 some reports4,5 of the modification of large bandgap micro- electrode, we adopted the solution growth technique described porous semiconductor particulate films using inorganic short in ref. 9, which is an attractive, low-cost and simple method bandgap quantum-sized semiconductor particles such as for large area deposition onto various substrates with various ultrafine CdS particles to embed into the microporous TiO2 materials as well as complex geometry.9,10 We immersed the electrode and to use these modified layers as light-converting microporous TiO2 electrode in the deposition bath composed electrodes. Visible light was absorbed by the ultrafine particles of 5 ml of 1 mol dm-3 cadmium acetate, 2.5 ml of 7 mol dm-3 which, consequently, transferred electrons into the TiO2 sub- triethanolamine, 5 ml of 13 mol dm-3 ammonia and 5 ml of strate.It has been shown by others6,7 that the photoresponse 1 mol dm-3 thiourea at 40°C. After about 6 h, the TiO2 of a large bandgap semiconductor can be extended into the electrode covered with an orange–yellow deposit was taken visible range by coupling with a short bandgap semiconductor.out, washed with distilled water and dried in air, to give the Since these ultrafine particles exhibit chemical and physical CdS modified electrode. properties which dier markedly from those of the bulk solid as well as of the individual molecules, very high quantum Preparation of dye-modified electrodes yields for the electron transfer process were obtained.In a previous paper,8 we described the preparation and The dye gallium tetrasulfophthalocyanine (GaTSPc) was syncharacterization of a TiO2 microporous electrode as well as thesized in our institute and its structure is shown in Fig. 1. the photovoltaic properties of such an electrode sensitized with Dye molecules were coated on the TiO2 electrode by immersing a dye.In the present work, a microporous titanium dioxide the electrode in a methanol solution with a dye concentration (TiO2) electrode has been prepared on a conducting glass support. Then we included Q-CdS particles and GaTSPc dye molecules, which possess dierent spectral absorption regions in the visible range, to co-modify the TiO2 electrode.Compared with electrodes using either Q-particles or organic dyes, a better spectral match to the solar spectrum was obtained for a microporous TiO2 electrode sensitized by a combination of Q-CdS particles and phthalocyanine dye molecules. Experimental Preparation of TiO2 electrode The nanocrystalline TiO2 films were fabricated as follows. A colloidal solution was prepared by adding commercial TiO2 power (P25, Degussa, a mixture of ca. 25% rutile and 75% anatase, surface area 55 m2 g-1, diameter 25 nm) to a small amount of water and surfactant (Triton x-100). The colloidal solution was then coated onto a conducting glass support. Fig. 1 The structure of GaTSPc, M=GaOH; R=SO3- After air drying, this was heated in air in a Lindberg tube J.Mater. Chem., 1997, 7(5), 737–740 737of about 10-4 mol dm-3. After at least 3 h of soaking, we combine to form CdS. The CdS formed homogeneously throughout the solution grows into clusters which can diuse obtained a dye-modified CdS/TiO2 electrode. to the substrate, where they aggregate to form a layer. The TiO2 electrode used as the substrate was microporous, and Experimental setup when the substrate was immersed in the solution, both The morphologies of the TiO2 and CdS/TiO2 electrodes were (NH2)2CS and Cd(TEA)2+ would penetrate into the TiO2 examined with an atomic force microscope (AFM) (DICO, electrode through the pores.Therefore, while some CdS par- NanoScope 3). The absorption spectra were recorded with a ticles were formed on the surface of the electrode, some CdS Shimadzu UV-2201 spectrometer.A two-electrode photoelec- particles were also deposited in the pores of the TiO2 electrode, trochemical (PEC) cell, composed of a modified TiO2 electrode, as was demonstrated in ref. 4. Probably owing to the spatial a counter electrode consisting of conducting glass sputter- confinement, the CdS particles inside the electrode were smaller coated with a 1 mm Pt film, and an electrolyte containing than those on the surface of the electrode.In order to estimate 0.1 mol dm-3 Na2S and 0.01 mol dm-3 Na2SO4, was used for the amount of CdS particles coated on the microporous photovoltaic studies. The cell had an area of 0.5 cm2. The TiO2 electrode, we allowed the CdS particles to dissolve in short-circuit photocurrent was measured with a Model CMBP- HNO3.After washing the electrode with water, a value of 1 potentiostat. Monochromatic illumination was obtained 361.8 mg cm-2 of Cd2+ (r=4.82 g cm-3) was found by atomic using a 500W xenon arc lamp in combination with a grating absorption spectroscopy. Considering the diameter of CdS monochromator, model WPG3D. The light intensity was particles as well as the roughness factor of the bare TiO2 calibrated using a model OM-1001C radiometer/photometer. electrode, approximately one or two layers of CdS particles were estimated to coat the TiO2 electrode.The interconnected ultrafine particles and pores on the TiO2 electrode form a Results and Discussion three-dimensional network structure. These pores not only AFM studies increase the surface area and allow the electrode to adsorb more monolayer dye molecules, but also enable the adsorbed AFM pictures of the surfaces of the TiO2 and CdS/TiO2 dye molecules to come into direct contact with the electrolyte.electrodes are shown in Fig. 2. The TiO2 film is microporous, When dye molecules absorb incident light, the photogenerated composed of interconnected particles and pores.It can also be electrons can transfer directly into the conduction band of seen that the TiO2 particles are uniform with an average semiconductor particles; at the same time, the holes remaining diameter of about 70 nm, which is much larger than the size in the dye can be rapidly removed by electron transfer from of the TiO2 particles in colloidal solution.Further experimental the redox system in the electrolyte, thus decreasing the prob- results show that the ating process is crucial to the morability of recombination and increasing the photoresponse. phology of the TiO2 electrode. The size of the particles in the TiO2 film increases with increasing temperature. The AFM Absorption characteristics and photovoltaic study study of the CdS/TiO2 surface indicates the formation of small particles on the surface of the TiO2 electrode, a trace section The absorption spectra of the electrode at various stages of analysis across the particles gives the size of these particles at modification are shown in Fig. 3.A conducting glass support about 10 nm in diameter. They are coated on the TiO2 surface is used as reference in the measurement.The TiO2 contains in the form of clusters. Considering the solution deposition 25% rutile and 75% anatase and the bandgap of rutile is 3 eV, process, this value of the particle size, however, is only the as compared to 3.2 eV for anatase, corresponding to funda- upper limit. mental absorption edges of 413 nm and 388 nm, respectively; The chemical reaction that leads to the formation of CdS is therefore the bare TiO2 film exhibits the fundamental absorp- as follows9 tion edge of rutile. Curve (b) is the absorption spectrum of the Q-CdS/TiO2 electrode.It shows that the CdS/TiO2 electrode Cd(TEA)2++(NH2)2CS+2OH- displaces the characteristic absorption of CdS well in to the �CdS+TEA+(NH2)2CO+H2O visible region with an apparent absorption shoulder around 500 nm and an weak absorption tail.The low-energy tail arises The chemical bath deposition mechanism is essentially based on the slow release of S2- ions from (NH2 )2CS and Cd2+ from the large particles, while the high-energy absorption shoulder arises predominantly from the small particles. The from Cd(TEA)2+ in the solution. The free Cd2+ and S2- Fig. 2 AFM pictures of the TiO2 (a) and CdS/TiO2 (b) electrodes 738 J. Mater. Chem., 1997, 7(5), 737–740be noted that both GaTSPc–CdS/TiO2 and GaTSPc/TiO2 electrodes show significant photocurrent responsescorresponding to the monomeric absorptionof GaTSPc. The photocurrent of the GaTSPc–CdS/TiO2 electrode is about three times that of GaTSPc/TiO2 at the excitation wavelength in the red region.The improvement of the sensitization eciency can be explained as follows. First, owing to the absence of an eective space charge layer at the particle/electrolyte interface,12 by utilizing nanocrystalline semiconductor particulate films in photoelectrochemical cells, the photogenerated charge carrier can move in both directions under this circumstance. Thus, the photogenerated electrons either recombine readily with holes or leak out at the electrolyte interface, instead of flowing through the external circuit.However, Q-CdS particles were coated onto the surface of the TiO2 particles to form a CdS/ Fig. 3 Absorption spectra of bare TiO2 (a), CdS/TiO2 (b), GaTSPc– TiO2 coupled structure. In this coupled system, the injected CdS/TiO2 electrodes and GaTSPc in methanol solution (d); all except charge from the excited state of the dye molecules quickly (d) used a conducting glass support migrates from CdS into the lower lying conduction band of TiO2; meanwhile, the junction between the TiO2 and CdS particles creates an energy barrier against the flow of electrons absorption of the GaTSPc–CdS/TiO2 electrode further extends towards the electrolyte, thus reducing the possibility of reverse the absorption range of the electrode into the red and displaces electron flow, i.e.the CdS/TiO2 coupled structure exhibits a the characteristic absorption of GaTSPc with a maximum in beneficial rectification eect, so an eective charge carrier the visible region around 610 nm. The shoulder around 690 nm separation and transport throughout the particulate film in curve (c) of Fig. 3 is assigned as the monomeric GaTSPc Q- becomes feasible. The recombination and electron leakage band.11 The broad band around 610 nm is enhanced compared losses are relatively low. In addition, a TiO2 electrode coated with the spectrum of GaTSPc in dilute methanol solution (d), with smaller Q-CdS particles possesses a larger surface area indicating the occurrence of aggregation11 of GaTSPc mol- than the bare TiO2 electrode. As a result, a relatively high ecules on the electrode.Compared with the CdS/TiO2 elec- photocurrent response is generated in this system. trode, GaTSPc–CdS/TiO2 displays ecient absorption over a wider spectral range. Higher solar energy conversion eciency may be expected, since more incident light was absorbed by the GaTSPc–CdS/TiO2 electrode.We can conclude from the Conclusion above that modification of TiO2 with Q-CdS particles extends the optical absorption spectrum into the visible region, and Based on the present results and the above considerations, this that the absorption spectrum of the CdS/TiO2 electrode can paper can be summarized as follows.(1) A microporous TiO2 be further extended into the red region by sensitization with electrode can be modified with Q-CdS particles by the chemical GaTSPc dye molecules. deposition method, and a CdS/TiO2 electrode can be further Fig. 4 shows the photocurrent action spectra of the CdS/ sensitized with GaTSPc molecules. (2) Compared with conven- TiO2 electrode (a), GaTSPc–CdS/TiO2 (b) and GaTSPc/TiO2 tional electrodes using either CdS particles or phthalocyanine, electrodes (c).The photocurrent action spectrum of the Q- modification of the TiO2 electrode with both CdS particles CdS/TiO2 electrode correlates well with its absorption spec- and GaTSPc not only results in a better spectral match to the trum, indicating that Q-CdS contributes significantly to the solar spectrum, but also extends the photoresponse of the photocurrent of the electrode.It is obvious that the photo- electrode to cover most of the visible range. (3) The monomeric current action spectra of the GaTSPc-CdS/TiO2 and GaTSPc/ molecules contribute to the spectral sensitization eect on the TiO2 electrodes are very dierent from the absorption spectra CdS/TiO2 electrode, while the contribution of the GaTSPc in that the maximum of the curve changes from about 610 nm molecules in aggregated form is negligible.(4) The rectification to 690 nm. We infer from the above that the monomeric of the coupled CdS/TiO2 system can eectively increase the GaTSPc molecules have contributed to the photocurrent of response of the electrode. (5) Among the phthalocyanines (such the electrode, while the GaTSPc molecules in the aggregated as GaTSPc, ZnTSPc, InTSPc, VoTSPc, etc.) studied in our form contribute little or nothing to the photocurrent.It should experiments, GaTSPc shows the best sensitization eect on the electrode. This work is supported by the Natural Science Foundation of China, project no. 69371032. References 1 B.O’Regan and M. Gra�tzel, Nature (L ondon), 1991, 253, 737. 2 B. O’Regan, J. Moser andM. Gra�tzel, J. Phys. Chem., 1990, 94, 872. 3 B. O’Regan, M. Gra�tzel and D. Fitzmaurice, Chem. Phys. L ett., 1991, 183, 89. 4 R. Vogel, K. Pohl and H.Weller, Chem. Phys. L ett., 1990, 174, 241. 5 R. Vogel, P. Hoyer and H.Weller, J. Phys. Chem., 1994, 98, 3183. 6 A. Ennaoui, S. Fiecher, Ch. Pettenkofer, N. Alsonso-Vante, Fig. 4 Photocurrent action spectra of incident photon conversion K. Buker, M. Bronold, Ch. Hopfnev and H. Trpbutsch, Sol. Energy Mater. Sol. Cells, 1993, 29, 289. eciency (IPCE) of CdS/TiO2 (a), GaTSPc–CdS/TiO2 (b) and GaTSPc/TiO2 (c) electrodes. Electrolyte 0.1 mol dm-3 Na2S and 7 A. Ennaoui, S. Fiechter, H. Tribatsch, M. Giersig, R. Vogel and H. Weller, J. Electrochem. Soc., 1992, 139, 2514. 0.01 mol dm-3 Na2SO4 J. Mater. Chem., 1997, 7(5), 737–740 7398 Y. C. Shen, L.Wang, Z. H. Lu, Y. Wei, Q. F. Zhou, H. F. Mao and 11 J. R. Darwent, P. Douglas, A. Harriman, G. Porter and M. C. Richoux, Coord. Chem. Rev., 1982, 44, 83. H. J. Xu, T hin Solid Films, 1995, 257, 144. 9 A. Mondal, T. K. Chaudhuri and P. Pramanic, Sol. Energy Mater., 12 A. Henglein, T op. Curr. Chem., 1988, 143, 115. 1983, 7, 431. 10 R. A. Boudred, D. K. Pandga and K. L. Chopra, J. Electrochem. Paper 6/07318B; Received 28th October, 1996 Soc., 1983, 130, 513 Mater. Chem., 1997, 7(5), 737–740

ISSN:0959-9428    

DOI:10.1039/a607318b

出版商:RSC    

年代:1997

数据来源: RSC

摘要:

Zn K-edge EXAFS study of SILAR-grown zinc sulfide thin films Seppo Lindroos,*a† Yves Charreire,b Tapio Kannianinen,a Markku Leskela�a and Simone Benazethc aDepartment of Chemistry, University of Helsinki, PO Box 55, FIN-00014 University of Helsinki, Finland bUniversite� Paris VI, 5 Place Jussieu, F-75252 Paris cedex 05, France cL .U.R.E. L aboratoire de Chimie Physique, Faculte� de Pharmacie, Universite� de Paris Sud, F-91405 Orsay, France Zinc sulfide thin films grown by the successive ionic layer adsorption and reaction (SILAR) method have been characterized by extended X-ray absorption fine structure (EXAFS) measurements. The ZnS films were well crystallized even as-grown but annealing improved the crystallinity clearly.The films contained small amounts of oxygen and, according to the EXAFS results, oxygen in the SILAR-grown zinc sulfide thin films occurred as hydroxide ions both in the as-grown and in the annealed samples. The simulated radial distribution curves for two models, Zn(OH)2/ZnS and ZnO/ZnS, were calculated to analyse the film composition.The ZnS thin films were also characterized by IR and electron spectroscopy for chemical analysis (ESCA) measurements.The successive ionic layer deposition and reaction (SILAR) Experimental method based on a heterogeneous reaction on the solid/liquid Sample preparation interface was introduced by Nicolau in 1985.1 The film growth in the SILAR method consists of four steps: in the first step The zinc sulfide thin films were grown by the SILAR method cations are adsorbed on the substrate surface, in the second with deposition equipment described earlier.2 The aqueous step all cations not adsorbed are rinsed away with purified precursor solutions used in the deposition were 0.1 M ZnCl2 water.The solvated anions enter the diusion layer in the next (pH=5.3) and 0.05 M Na2S (pH=11.8). The immersion time reaction step and they react with the adsorbed cations.The was 20 s and the rinsing time was 100 s. The flow rate of the ions which have not reacted are again washed away with a purified rinsing water was 600 ml min-1. The ZnS thin films rinsing pulse in the fourth step. After a complete cycle a solid were grown on indium tin oxide (ITO)-covered glass substrates. compound is formed on the surface of the substrate. By During the growth the substrate was kept in an N2 environ- repeating these steps a thin film can be grown layer by layer.ment. Sample A was as-grown and sample B was annealed for Hence the thickness of the film can be controlled directly by 3 h at 500 °C in a nitrogen atmosphere. Both samples were the number of deposition cycles. deposited simultaneously and had thicknesses of 164 and The main advantage of the SILAR method is the possibility 150 nm respectively.Chemical and physical characterization to deposit thin films that are impossible or dicult to grow (chemical analysis, UV, XRD, RBS, SEM) of the SILAR- using gas-phase methods. On the other hand, due to the grown zinc sulfide thin films have been published earlier.2 layerwise growth of the thin film, the quality of the SILAR- For comparison to the SILAR-grown ZnS thin films, ZnO grown films is comparable with films grown with other liquid- and ZnS powders, amorphous ZnS thin films prepared by phase thin film deposition methods such as chemical bath sputtering and well crystallized thin film ZnS samples made deposition.1–7 The SILAR technique has so far been used to by the ALE (atomic layer epitaxy) method in the Laboratory grow II–VI compounds, such as ZnS.1,2 The low temperature of Inorganic and Analytical Chemistry at the Helsinki of the SILAR deposition allows also the use of temperature- University of Technology were also studied.The hexagonal sensitive precursors and substrates.8 forms of ZnS and ZnO were used. One of the main problems in the SILAR process has been that some oxygen remains in the films owing to the aqueous EXAFS measurements growth conditions as confirmed by Rutherford backscattering spectroscopy (RBS) measurements.2 According to the RBS EXAFS experiments were performed at Lure (Orsay, France) results oxygen in the ZnS thin films is partly in water and using X-ray synchrotron radiation emitted by the DCI storage partly bound to zinc as hydroxide or oxide.The form of ring (1.85 GeV, 350 mA), in the EXAFS II installation. Data oxygen, viz. oxide or hydroxide ion, in ZnS films is an were collected using an Si(111) double crystal monochromator. interesting and important question when considering possible Rejection of harmonics was not adopted. The EXAFS spectra applications, and cannot be answered by Rutherford backscatt- of the thin films were recorded in a total electron yield ering measurements only.mode9 and measurements were made around the zinc K edge This study was carried out on one hand to analyse the (9659 eV). oxygen concentration and coordination in the SILAR-grown The EXAFS data analysis was carried out using a PC films, i.e. whether the oxygen is bound to zinc as oxide or program written by Bonnin et al.10 The background absorption hydroxide.On the other hand, the purpose of this study was features which are superimposed on the EXAFS oscillations to compare the crystallinity of the SILAR-grown ZnS thin were removed by linear extrapolation of the pre-edge. The films to films grown by sputtering or gas-phase atomic layer normalized EXAFS signal x(E)=[m(E)-m(E0)]/m(E0) was con- epitaxy (ALE) methods.verted to wavevector k space, using k=2m(E-E0)1/2/B, where The SILAR-grown ZnS thin films were also characterized E0 is the threshold energy close to the absorption edge of Zn by IR spectroscopy and ESCA measurements. (maximum of the first derivation of the spectrum).11 For the x(k) simulation, DE0 is a parameter corresponding to a shift of the experimental value of E0.Fourier transformation was performed using a Kaiser window in the range 2<k<13 A° , † Email: SLindroos@phcu.Helsinki.Fi J. Mater. Chem., 1997, 7(5), 741–745 741with a k3 weighting.12 After this the EXAFS contribution of IR and ESCA measurements the chosen coordination sphere was isolated by Fourier filtering The IR spectra were recorded using a Bruker IFS 85 spec- and the experimental x(k) was obtained by back-Fourier trans- trometer and for powder samples the CsBr pellet technique form.After the back-Fourier transform to the k space the was used. filtered EXAFS signal was k3 weighted and Fourier transformed ESCA measurements were carried out with a Phi small spot through a Kaiser window (t=3) on the spectral range ESCA 5400 electron spectrometer using unmonochromated extending from k=4 to ca. 12 A° -1. Fitting was performed Mg-Ka radiation. using the theoretical amplitude and phase functions of McKale et al.,13 or experimental functions obtained by pure ZnS and Results and Discussion ZnO reference compounds. The inelastic losses are usually approximated by the factor exp(-2rj/l) where l is the inelastic The structural information used in this study is presented in mean free path.To take into account the k-dependence of this Table 1 and the parameters used in EXAFS simulation and damping term we used the factor exp(-rjL /k) where L=2k/l. fits are presented in Table 2. Fig. 1 shows the x(k) oscillation The relative mean squared static and thermal disorder between of the sample A and for comparison the x(k) oscillation of a the central atom and the atoms in the jth coordination sphere ZnS thin film sample deposited by the ALE technique.Fig. 2 is given by sj2. To take into account this eect we introduced shows the k3x(k) Fourier transform [radial distribution func- the factor exp(-sj2k2). The Fourier transform was made with tion P(r)] in k space in the range 2–13 A° -1 for the SILAR respect to exp[2ikrj+wj(k)] where a phase shift wj(k) is needed samples.The same curves for ZnS thin films prepared by the to take into account the potentials due to both the central ALE method and by sputtering are presented for comparison. atom and the backscatterers. The phase shift is not totally The peak located at 2.00 A° (shift of -0.34 A° from the correct simulated in the prograis causes a shift of all the distance) is at the exact position of sulfur atoms in the ZnS peaks in P(r) to closer to the origin by aj=rj-r¾j, where r¾j is lattice.The first peak at 1.5 A° is due to oxygen and confirms the observed distance in P(r). From reference compounds ZnO thus the presence of oxygen in SILAR-grown films.The and ZnS the shift was estimated to be 0.34 A° for oxygen and intensity of the second (ZnMS) peak of samples A and B in sulfur peaks. The parameters (L , s) and (r, DE0) are not Fig. 2 is higher compared with that found in samples deposited independent. From earlier studies12,14 on ZnS we had a good by sputtering indicatinga better crystallinity in SILAR samples.estimation for the values of the parameters at the beginning On the other hand, ZnS thin films deposited by ALE from the of the fit. gas phase at elevated temperature and reduced pressure, are even better crystallized than SILAR samples. Annealing clearly approves the crystallinity of the SILAR-grown ZnS films. Table 1 Coordination shells (r/A° ) around Zn in ZnS, ZnO and Fig. 3 depicts the experimental radial distribution functions e-Zn(OH)2 of ZnO and ZnS reference compounds (powder samples were studied by the transmission method) and simulated radial shell ZnS ZnO Zn(OH)2 distribution function of the first six coordination spheres of e- Zn(OH)2. The simulation for Zn(OH)2 was used because both 1 2.34 (4 S) 1.99 (4 O) 1.95 (4 O) 2 3.82 (12 Zn) 3.23 (12 Zn) 3.43 (4 Zn) commercial and several synthesized zinc hydroxide samples 3.91 (1 S) 3.25 (1 O) were found to be mixtures of Zn(OH)2 and ZnO.The surround- 3 4.48 (9 S) 3.80 (9 O) 3.6–4.0 (4 O) ings of zinc atoms in the ZnO and Zn(OH)2 structures for the 4 5.41 (6 S) 4.57 (6 Zn) 4.05–4.7 (8 O) first neighbours (four oxygen atoms approximately at the same 5 5.47 (6 S) 4.60 (6 O) 4.8–4.9 (4 Zn) distance of 1.9 A° ) are similar and were observed from this Table 2 Bond lengths (r), coordination numbers (N), Debye–Waller factors (s), L factors, edge shifts (DE0) and standard deviations obtained by fitting EXAFS spectra for ZnS thin films (underlined values are fixed) compound shell N r/A° s/A° L /A°-2 DE0/eV .(Dx2/x2) fitting used 1–O 0.75 1.91 0.071 1.50 — 28 experimental (19±2 ZnS sample A 2–S 3.25 2.23 0.081 1.75 — atom% oxygen) ZnS sample B 1–O 1.15 1.99 0.064 1.50 — 29 experimental (29±3 (annealed) 2–S 2.85 2.30 0.068 1.75 — atom% oxygen) ZnS reference (ALE- 1–S 4 2.33 0.072 1.75 -4.7 26 McKale et al.13 used grown thin film) 2–Zn 12 3.83 0.115 1.75 -5.1 for Zn–S 3–S 9 4.46 0.114 1.75 -1.1 ZnO reference 1–O 4 1.90 0.069 1.50 -0.3 24 McKale et al.13 used (powder) 2–Zn 12 3.23 0.093 1.50 0 for Zn–O and 3–O 9 3.80 0.093 1.50 0 Zn–Zn Zn(OH)2 reference 1–O 4 1.95 0.069 1.50 — — experimental 2–Zn 4 3.43 0.093 1.50 — 3–O 4 3.80 0.093 1.50 — 4–O 4 4.20 0.093 1.50 — 5–Zn 4 4.90 0.093 1.50 — 6–O 4 4.60 0.093 1.50 — ZnO/ZnS 3–Zn 3.45 3.23 0.14 1.50 -7 — McKale et al.13 4–Zn 8.55 3.82 0.14 1.75 -7 5–O 2.59 3.80 0.14 1.50 -7 6–S 6.40 4.48 0.14 1.75 -7 Zn(OH)2/ZnS 1–O 1.15 1.92 0.064 1.50 -7 29 McKale et al.13 2–S 2.85 2.31 0.068 1.75 -7 3–Zn 1.25 3.43 0.14 1.50 -7 4–Zn 8.55 3.83 0.14 1.75 -7 5–O 2.30 4.00 0.14 1.50 -7 6–S 6.40 4.48 0.14 1.75 -7 742 J. Mater.Chem., 1997, 7(5), 741–745Fig. 1 A comparison of EXAFS oscillations x(k) for SILAR-grown ZnS thin film sample A (,) and ALE-grown ZnS thin film reference Fig. 3 Radial distribution function of ZnS (C) and ZnO (,) pow- (—) (prepared by the ALE group, Helsinki University of Technology) ders (transmission mode) and simulated P(R) curve for Zn(OH)2 (—) Fig. 2 Radial distribution curves P(R) for SILAR thin films samples Fig. 4 The P(R) curves for the ZnO/ZnS mixtures, where the percent- A (C) and B (—), ZnS film made by sputtering (A) and ZnS age of ZnO increases from 0 to 1, 2, 4, 8, 16, and 64 atom% reference thin film prepared by ALE ( ) Fourier transform (1.6 A° , shift-0.34 A° in Fig. 3). The ZnMOH (step in k space=0.05 A° -1, in R space=0.061 A° ) with the program used leading to poor spatial resolution and conse- bonding in the SILAR samples can be identified by the second coordination sphere.The peak of the four Zn atoms at 3.43 A°quently low accuracy for determination of oxygen. In this study the O/S ratio was first determined by the fit in Zn(OH)2 structure can be easily separated from the peak of twelve Zn atoms at 3.23 A° in ZnO and 3.82 A° in ZnS technique from the oxygen and sulfur peaks which correspond to a coordination number of four for the Zn atoms.In the structures. To gain an idea about the sensitivity of EXAFS determi- second step the experimental P(R) curves of the two thin films were compared with theoretical simulations made for two nation for the oxygen concentration the P(R) curves for the ZnO/ZnS mixtures, where the percentage of ZnO increases mixtures (ZnO)x/(ZnS)y and [Zn(OH)2]x/(ZnS)y where x and y are determined from the O/S ratio.By simulation and fitting from 0 to 1, 2, 4, 8, 16, and 64 atom% were studied (Fig. 4). Unfortunately there is a shoulder in the P(R) curve of ZnS techniques the oxygen content can be estimated: the as-grown SILAR sample A contained 19±2 atom% and the annealed close to the oxygen peak at 1.5 A° (shift -0.34 A° ). Almost 8 atom% of oxygen is required to observe a notable change in sample B 29±3 atom% oxygen.Annealing increases the EXAFS signal and increases the long-range order. The higher this shoulder. The oxygen peak is described only by 8 points J. Mater. Chem., 1997, 7(5), 741–745 743Fig. 6 Experimental (—) and simulated (,) EXAFS oscillations x(k) of the first peak (oxygen) and the second peak (sulfur) of sample B Fig. 5 The simulated radial distribution curve P(R) for the first peaks (O, S,Zn, Oand Zn) of the two models: mixed structures Zn(OH)2/ZnS (A) and ZnO/ZnS (,) with the experimental P(R) for the sample B imental and simulated x(k) curves for the isolated two first (—) peaks of sample B and the fit was very good. Fourier filtering between 2.39 and 4.23 A° gives the x(k) oscillations of the next four peaks in the ZnO/ZnS mixture oxygen concentration after annealing is probably due to the increased surface roughness of the ZnS thin films.This enables model and four or five peaks in the Zn(OH)2/ZnS mixture model (see Table 2). Fig. 7(a) and (b) show simulated x(k) and the chemisorption or physisorption of atmospheric oxygen at the surface of zinc sulfide and hence an oxidized layer is formed P(R) curves for these two models and for sample B.The simulation could not be made by the fit technique because on it. Also it has been observed, according to optical absorption studies, that the annealing of ZnS thin films can decrease the there are too many free parameters but s increased to 1.4 A° between the third and the sixth coordination spheres of the thickness of the film due to partial sublimation of ZnS even at low temperatures.The increase of oxygen content in the model. The simulations of the P(R) and x(k) curves are good enough to conclude that the SILAR films consist of a annealed SILAR-grown ZnS thin film can be attributed to the selective sublimation of ZnS. ZnS/Zn(OH)2 mixture. According to the XRD measurements, the ZnS and ZnS5Mn The simulated radial distribution curves P(R) for the first peaks (O, S, Zn, O and Zn) of the two models, mixed structures thin films prepared by SILAR were polycrystalline.2,5 Annealing of the ZnS5Mn thin film at 500°C in an N2 Zn(OH)2/ZnS and ZnO/ZnS, with the experimental P(R) for the sample B are presented in Fig. 5. It can be concluded from atmosphere slightly increased the intensity of the [111] peak at 2h=28.5° and decreased the FWHM.5 The oxygen content the curves that the ZnO/ZnS model must be incorrect.The model Zn(OH)2/ZnS resembles more closely the experimental of the SILAR-grown ZnS films, determined by RBS, varies from 10 to 20 atom%.2 In the ZnS5Mn thin films the oxygen P(R) but dierences in intensity can be seen at larger distances. Annealing increased both the EXAFS signal and the long- content decreased remarkably in the annealing process from 20 to 10 atom%.5 The hydrogen content, determined by range order of the structure.Also, the resolution between oxygen and sulfur was better after annealing with the error for nuclear reaction analysis (NRA), of the ZnS5Mn thin films also decreased from about 10 atom% of the as-deposited film the determination of the O/S ratio being probably 30%.The errors indicated in Table 2 concerning the oxygen proportion to 0.2–0.4 atom% of the annealed film.5 The dierence between the EXAFS and NRA results is possibly because the infor- correspond only to the interval explored during the fit. Nevertheless, we can conclude that both samples contain mation depth of EXAFS is smaller compared with the information depth of NRA and RBS.On the surface of ZnS5Mn relatively significant amounts of oxygen. However the intensity of the EXAFS peak increases as the crystallinity of the sample thin films there is a 10–20 nm thick layer where the hydrogen content is up to 2.5 atom% in the annealed films and even increases. Owing to the lower crystallinity in sample A the EXAFS results are more uncertain than in sample B.On the higher (ca. 5 atom%) in the as-grown films.5 Oxygen in this surface layer exists most probably as water and zinc hydroxide. other hand the better crystallinity in sample B may lead to a slight overestimate of the oxygen content. In the ZnS thin films grown by a chemical deposition method the films contained significant amounts of oxygen either as Using the Fourier filtering technique, where one peak can be isolated in the experimental P(R) curve, the corresponding oxide (ZnS0.5O0.5) or as hydroxide [ZnS0.5(OH)].15 According to the ESCA measurements of ZnS thin films coordination sphere can be analysed separately.The two first peaks (O, S) in the range 1.29–2.39 A° were isolated and the grown by SILAR the Zn5S ratio was 151 and on an as-grown surface oxygen was detected.The oxygen content of the surface back Fourier transform of the selected P(R) gave the x(k) contribution of these two first peaks. The resulting ‘filtered’ decreased remarkably after the sample was sputtered three times. It has been found that there is a thin oxidized layer on EXAFS spectrum [k3x(k)] was then fitted. The result was a model where the coordination number is four and in the the surface of colloidal ZnS.The oxygen content of this surface layer diminished rapidly with depth and at a depth of 10 nm coordination sphere 19 and 29% of the atoms were oxygen in the A and B samples, respectively. Fig. 6 presents the exper- the oxygen signal halved in magnitude.16 744 J.Mater. Chem., 1997, 7(5), 741–745ZnMO(HCO2) by Ga� rd et al.20 was observed in the spectra. As a conclusion of the IR measurements the SILAR-grown ZnS thin films contain oxygen as zinc oxide and carbonate and water was also found in the films. The IR spectrum can be used only qualitatively to analyse oxygen since addition of only 1% ZnO to ZnS powder has been known to give a strong absorption at 460 cm-1.21 Conclusions The oxygen content and the form of the oxygen ion in the SILAR-grown zinc sulfide thin films has been determined by EXAFS measurements and simulation of ZnO/ZnS and Zn(OH)2/ZnS mixture models.The ZnS thin films prepared by SILAR contain 10–30% oxygen. Near the surface of the film, in the volume where EXAFS information is originated, the oxygen is bound more probably as ZnMOH than as ZnMO both in as-grown and annealed films.The absence of the peak corresponding to the twelve ZnMZn distances at 3.23 A° , characteristic of ZnO, confirms this conclusion. The crystallinity of the SILAR-grownZnS thin films is fully compatible with the ZnS thin films grown by the ALE method at higher temperatures.We thank Prof. J. Va� yrynen (University of Turku) and Mrs. H. Lepesant (Ecole National Superieure de Chimie de Paris) for ESCA and IR measurements, respectively. The research has been supported by Academy of Finland and Technology Advancement Centre of Finland (TEKES). References 1 Y. F. Nicolau, Appl. Surf. Sci., 1985, 22/23, 1061. 2 S. Lindroos, T. Kanniainen and M. Leskela�, Appl.Surf. Sci., 1994, 75, 70. 3 Y. F. Nicolau and J. C. Menard, J. Cryst. Growth, 1988, 92, 128. 4 Y. F. Nicolau and J. C. Menard, J. Appl. Electrochem., 1990, 204, 1063. 5 S. Lindroos, T. Kanniainen,M. Leskela� and E. Rauhala, T hin Solid Films, 1995, 263, 79. 6 R. Ortega Borges, D. Lincot and J. Vedel, 11th E.C. Photovoltaic Solar Energy Conference,Montreux, 1992, p. 862. 7 J. M. Don�a and J. Herrero, J. Electrochem. Soc., 1994, 141, 205. 8 S. Lindroos, T. Kanniainen and M. Leskela�, J.Mater. Chem., 1996, 6, 1497. 9 G. Tourillon, E. Dartyge, A. Fontaine, M. Lemonnier and F. Bartol, Phys. L ett. A, 1987, 121, 251. 10 D. Bonnin, P. Kaiser, C. Fretigny and J. Desbarres, Structures Fines d’Adsorption des Rayons X en Chimie, Ecole du C.N.R.S., Orsay, 1989, vol. 3, part 2. 11 H. D. Abrun�a, J. H. White, M. J. Albarelli, G. M. Bommarita, M. J. Bedzyk and M. McMillan, J. Phys. Chem., 1988, 92, 7045. Fig. 7 (a) Comparison between the back Fourier transform of exper- 12 Y. Charreire, A. Marbeuf, G. Tourillon, M. Leskela�, L. Niinisto�, imental radial distribution function and the x(k) simulation of the E. Nyka�nen, P. Soininen and O. Tolonen, J.Electrochem. Soc., third, fourth, fifth and sixth shells for the two mixture models 1992, 139, 619. Zn(OH)2/ZnS and ZnO/ZnS: (—) back Fourier transform of filtered 13 A. G. McKale, B. W. Veal, A. P. Paulikas, S.-K. Chan and P(R) of sample B; (C) simulation of a priori ZnO/ZnS model; (,) G. S. Knapp, J. Am. Chem. Soc., 1988, 110, 3763. simulation of a priori Zn(OH)2/ZnS model. (b) Radial distribution 14 Y.Charreire, D.-R. Svoronos, I. Ascone, O. Tolonen, L. Niinisto� function P(R) of EXAFS oscillations x(k) of Fig. 7(a) for k space: (—) and M. Leskela�, J. Electrochem. Soc., 1993, 140, 2015. the experimental curve; (C) a priori ZnO/ZnS model; (,) a priori 15 B. Mokili, M. Froment and D. Lincot, J. Phys. IV, 1995, 5, C3. Zn(OH)2/ZnS model. 16 D. T. Atkins and R. M. Pashley, L angmuir, 1993, 9, 2232. 17 K. Atherton, G. Newbold and J. A. Hockey, Discuss. Faraday Soc., In the IR spectra of the ZnS thin films deposited by SILAR 1971, 52, 33. the ZnMS (zinc sulfide) and ZnMO (zinc oxide) stretch 18 W. Hertl, L angmuir, 1988, 4, 594. 19 J. W. Kauman, R. H. Hauge and J. L. Margrave, J. Phys. Chem., vibration peaks were found at 302.8 and 468.6 cm-1, respect- 1985, 89, 3541. ively. The broad moisture peak at 3416 cm-1 possibly masks 20 R. Ga�rd, Z-X. Sun and W. Forsling, J. Colloid Interface Sci., 1995, two narrow peaks at ca. 3670 and 3620 cm-1 and two broad 169, 393. peaks at ca. 3555 and 3440 cm-1 of hydroxy groups.17,18 The 21 Q. Yitai, S. Yi, C. Qianwang and C. Zuyao, Mater. Res. Bull., 1995, water bending peak at 1646.1 cm-1 was also identified. The 30, 601. ZnMOH stretching peak at 648 cm-1 19 could not be observed. The broad absorption band at 1457–1377 cm-1 assigned as Paper 6/08383; Received 13th December, 1996 J. Mater. Chem., 199

ISSN:0959-9428    

DOI:10.1039/a608383h

出版商:RSC    

年代:1997

数据来源: RSC

摘要:

Giant magnetoresistive thin films of (La,Pr)0.7(Ca,Sr)0.3MnO3 prepared by aerosol MOCVD O. Yu. Gorbenko,*a A. R. Kaul,a N. A. Babushkinab and L. M. Belova,b aChemistry Department,Moscow State University, 119899 Moscow, Russia bKurchatov Institute, 123182Moscow, Russia It has been demonstrated that high quality thin films of Ln1-xMxMnO3 can be successfully prepared by aerosol MOCVD at 750 °C from volatile thd complexes.Subsequent annealing in oxygen at 750°C is necessary to stabilize the oxygen content of the films. XRD patterns of the films showed them to be pseudocubic with an apparent lattice parameter linear in the average ionic radius of Ln and M. The evolution of Ln1-xMxMnO3 film morphology with increasing the film thickness has also been studied. The morphological instability of the MOCVD process results in the formation of a hillocky surface with thickness >2000 A° .The electrical properties of La1-xCaxMnO3 and La1-xSrxMnO3 correlate with those reported for bulk and thin film materials. The substitution of Pr for La in La1-xSrxMnO3 reduces the maximum resistivity temperature, Tp, non-linearly. La1-xCaxMnO3 thin films reveal a shift of Tp downwards in the case of substrate materials with a positive lattice mismatch with the films.La0.35Pr0.35Ca0.3MnO3/LaAlO3 demonstrates a very complex temperature dependence of the resistivity, which is described using a conceptual phase diagram of Ln1-xMxMnO3. A marked GMR eect was observed for La0.35Pr0.35Ca0.3MnO3/LaAlO3 below 21 K (ca. 1010%) and at ca. 70 K even in a field of 1 T.Ln1-xMxMnO3 perovskites, where Ln3+ is a rare-earth-metal (Ln=La, Pr and M=Sr,Ca) thin films by aerosol MOCVD and studied their structures and electrical properties. cation and M is a doubly charged cation with a large ionic radius, with both Ln and M populating the A positions of the ABO3 perovskite lattice, have attracted considerable attention Experimental owing to the recent discovery of giant negative magneto- Aerosol MOCVD, which was used to prepare the thin films, resistance (GMR) of such compounds.1 The GMR ratio r= involves deposition from precursor vapours produced by the [R(0)-R(H)]/R(H), which can be easily reduced to R(0)/R(H), evaporation of their organic solutions nebulized in the carrier where R(0) and R(H) represent the resistance in zero field and gas flow.22,23 This method was found to combine reprodu- that in an applied magnetic field H=1–10 T (irrespective of cibility of vapour composition with flexibility of composition whether H is the saturation field or not), characterizes the variation.Deposition conditions are summarized in Table 1. eect. A variety of possible applications2 such as sensors, The Ln0.7M0.3MnO3 thin films prepared are listed in Table 2.reading heads, random access memory and hard disks with Volatile precursors were as follows: La(thd)3, Pr(thd)3, areal densities up to 10 Gbit in-2, requires not only a large r Sr(thd)2, Ca(thd)2 and Mn(thd)3 (thd=2,2,6,6-tetramethyl- but also a high sensitivity to the driving field dr(H)/dH as heptane-3,5-dionate).All precursors were synthesized by well as a relatively low saturation field. The most widely standard techniques and sublimed in a vacuum before use. studied perovskites, La1-xMxMnO3,3–9 leave much to be desired to satisfy these demands. That the role of the average Table 1 Parameters of the deposition process ionic radii of the cations in the A positions (rA) is crucial was recently explained.4,10,11 By decreasing rA, a higher r can be parameter value achieved.The nature of the substituting cation is also important. Changing the M cation to a cation with a smaller rA stagnation flow reactor type cold wall shifts the temperature of the second-order phase transition deposition temperature/°C 750 total pressure/mbar 6 (P�FMM) from a high-temperature paramagnetic phase (P) total gas flow/l h-1 20 with activated conductivity to a low-temperature ferromagnetic partial pressure of O2 /mbar 3 phase with metal-like conductivity (FMM) downwards.10,12 deposition rate/mm h-1 0.5–1 The variation of Ln has a similar but more complicated eect.solventa consumption/ml h-1 20 A variety of new phase states appears for Ln=Pr or Nd.13–16 evaporator temperature/°C 250 Finally, the x range where a metal-like ferromagnetic phase aDiglyme (CH3OCH2CH2OCH2CH2OCH3) with 5 vol.% Hthd exists is restricted to 0.2<x<0.5 with a dome-shaped dependence of the magnetization on x.17 Thus, it is reasonable to Table 2 Thin films of (La1-x,Prx )0.7M0.3MnO3 prepared in this work stay in the middle of this range in order to analyse the eect of Ln or M variation correctly.10 M x abbreviation substratesa Comparable values of r were observed for both ceramics and thin films, causing us to reject the initial hypothesis which Sr 0 LSM MgO, YSZb, LaAlO3 had connected high r values with thin epitaxial films.18,19 Ca 0 LCM MgO, YSZ, LaAlO3, SrTiO3 Sr 11xf.0 PSM MgO, LaAlO3 Nevertheless, the electrophysical properties reported for films Ca 1 PCM MgO, LaAlO3 and bulk materials of practically the same composition can Sr 0.5 LPSM MgO, LaAlO3 dier drastically (compare for instance ref. 10 and 20), and Ca 0.5 LPCM MgO, YSZ, LaAlO3 even the substrate material can influence the phase transitions in the manganite films.21 aSubstrate orientation (001) in the pseudocubic system. bYSZ= ZrO2(Y2O3) In the work reported here we prepared Ln0.7M0.3MnO3 J.Mater. Chem., 1997, 7(5), 747–752 747XRD [four-circle diractometer, Siemens D5000 with oxygen stoichiometry determination for the Ln1-xMxMnO3+d films has not yet been described. Even in the case of the bulk secondary graphite monochromator (Cu-Ka radiation)] was applied to determine the phase composition, orientation and materials there is a lack of data.Reasons for this include the absence of a definite reference state (three oxidation states of lattice parameters of the prepared films. h–2h, Q scans and rocking-curve measurements were used. SEM was accom- manganese, namely Mn4+, Mn3+, Mn2+, occur in the material with the change of oxygen stoichiometry24,25), sample pre- plished by CAMSCAN equipped with an EDAX system for quantitive analysis.SNMS (sputtered neutrals mass spectro- history eects owing to cation disorder with a high density of vacancies in both A and B cation sublattices and an oxygen metry) depth profiling was carried out with an INA-3 system. Electrical resistivity measurements were carried out in a con- diusion rate which is too slow to obtain reliable results for bulk samples at low temperature. ventional four-point probe configuration using 3×10 mm2 bars cut from the 10×10 mm2 specimens.LaMnO3+d is capable of reduction with decreasing the oxygen partial pressure. According to ref. 26, at 900 °C and pO2=10-6 atm La1-xSrxMnO3+d has d#0 for 0x0.5. At Results and Discussion pO2=0.21 atm, d#0.15 was reported for LaMnO3+d at 730°C25 and d# 0.14 at 1100°C.26 With increasing x the d Post-deposition treatment range between the oxygenated and reduced forms decreased Post-deposition annealing is known to influence the electrical and tended towards 0 for x0.3.26 In fact, we have observed properties of Ln1-xMxMnO3.5,19 For the preliminary experi- a less profound change of the XRD pattern after annealing an ments we prepared La1-xSrxMnO3+d films with x=0.15 and La0.7Sr0.3MnO3+d film as compared to La0.85Sr0.15MnO3+d. 0.30. Various (La,Ln)0.7M0.3MnO3 samples prepared in ref. 10 As-grown films showed broad and sometimes split peaks of showed d#0 over a wide range of oxygen partial pressure. the perovskite phase in the XRD patterns (Fig. 1) with poorly Thus, we can expect that after the annealing our thin films reproducible lattice constants.It was suggested that such have nearly the same oxygen stoichiometry and this factor behaviour originated from incomplete oxidation at the cooling does not contribute to the dierences in their properties. stage. Since the deposition was carried out at low oxygen Consequently we omit d in the compound formulae written partial pressure (Table 1), and an additional depletion was below.possible because of the consumption ooxygen to oxidize solvent molecules on the film surface, the as-grown film was XRD characterization expected to be reduced. The fact that after the deposition, reoxidation started but was incomplete under the cooling XRD patterns of the annealed films are similar to those of a cubic lattice without any splitting of the solely observed (00l) conditions, resulted in the XRD pattern observed.The solution to the problem was found to consist of sub- or (hh0) reflections (in the pseudocubic system) or superstructural peaks. This does not mean that the perovskite film is sequent annealing of the films prepared by aerosol MOCVD in oxygen at 750 °C for 0.5 h. The choice of the annealing necessarily cubic because of the film orientation.The XRD pattern of the oriented perovskite film does not provide conditions was based on the data of van Roosmalen et al.,24 who had showed by TG that oxygen uptake by the LaMnO3 unequivocal data on the lattice symmetry or a full set of lattice parameters. If the parameters calculated from the (00l) or (hh0) lattice is negligible below 700 °C. After the treatment the diraction patterns changed significantly (Fig. 1) and became zones coincide, which was the case for our films, then we have an indication that the lattice is cubic. Apparent cubic lattice reproducible; simultaneously the room-temperature resistivity decreased to approximately half its previous value. The results parameters a for the films are shown in Fig. 2, which reveals a nearly linear dependence on the average ionic radius, rA.did not change after extra annealing under the same conditions. Consequently, the oxygen content of the film reached an The dependence is not trivial, because variation of rA results in tilting and rotation of the MnO6 octahedra building a three- equilibrium with the oxygen atmosphere at 1 bar and 750 °C, which provides the same reference state for the dierent film dimensional array rather than a simple contraction of bond distances, as shown for bulk samples.10,12 For the range of rA compositions under investigation.Thus, we used such postdeposition annealing for all compositions prepared. under consideration, orthorhombic or rhombohedral (at the lower temperature) phases in the bulk were described.A The d values were not determined in our work. The correct distortion of the array formed by MnO6 octahedra breaks up Fig. 2 Apparent lattice parameter, a, of (La,Pr)0.7(Sr,Ca)0.3MnO3 per- Fig. 1 Eect of subsequent annealing in oxygen at 750 °C for an ovskite films on MgO as a function of average ionic radius, rA , of the A site. The (002), (003), (110) and (220) peaks were used for the La0.85Sr0.15MnO3 film on YSZ: (a) (110) reflection, (b) (220) reflection.Dashed lines indicate the as-grown film, solid lines were obtained calculation. The a values found independently from (00l ) and (hh0) peaks dier by less than 0.006 A° for all samples (Table 2). after annealing. 748 J. Mater. Chem., 1997, 7(5), 747–752Fig. 3 Typical h–2h scans: (a) La0.7Sr0.3MnO3/LaAlO3, (b) La0.7Ca0.3MnO3/SrTiO3 the exchange of electrons governing magnetic properties and conductivity.Unfortunately, there is a lack of low-temperature XRD studies for the manganite thin films. Perovskite substrates (LaAlO3 a=3.793 A° , SrTiO3 a= 3.905 A° ) permit the deposition of layers showing only (00l ) reflections (Fig. 3); Q-scans prove the epitaxy of the ‘cube-oncube’ type of deposition.Rocking curves measured for the (00l ) reflections have FWHM=0.17–0.25° [Fig. 4(b)], which indicates the high epitaxial quality of the films. XRD patterns of Ln1-xMxMnO3/MgO reveal the (110) peak at thickness d>2000 A° . The large lattice mismatch between MgO (a=4.21 A° ) and Ln1-xMxMnO3 hinders the epitaxy of perovskites with (001)filmd(001)MgO.Deviation of Fig. 4 Representative rocking curves for the (002) reflections of the molar ratio Ln+M/Mn from unity results in the suppres- (La1-x,Prx)0.7M0.3MnO3: (a) La0.35Pr0.35Ca0.3MnO3 /MgO, sion of the (00l) orientation and the appearance of (hh0) peaks. (b) La0.7Ca0.3MnO3/SrTiO3 Analysis of the Q-scans shows that (00l)-oriented domains in the film on MgO have very good in-plane alignment: fourfold degeneracy of peaks is observed at a tilt angle equal to that between the plane under study and the (001) plane.Rocking curves for (00l ) planes with FWHM#0.5° were measured for the best Ln1-xMxMnO3/MgO films [Fig. 4(a)]. Ln1-xMxMnO3 on YSZ has a preferable orientation (110)fd(001)YSZ. Even a pure (110)fd(001)YSZ orientation can be obtained for the manganites.The result correlates to the diagonal match with YSZ (aÓ3 to 5.14×Ó2). The FWHM of the rocking curves of the (220) reflection was 0.58° for La0.7Sr0.3MnO3 on YSZ. Surface morphology Nearly stoichiometric films grow very smoothly up to until d<2000 A° . SEM images showed the featureless surface of such a film at a magnification of 5000–10000×. Small local disturbances were observed with increasing film thickness.In the range d=2000–3000 A° the surface changes from flat to hillocky. The hilliness becomes more marked for d>3000 A° , as was Fig. 5 AFM image of a 3500 A° thick La0.7Ca0.3MnO3 film on LaAlO3 . demonstrated also by AFM (Fig. 5). Such behaviour reflects The average roughness is 140 A° . the well known morphological instability of the CVD process controlled by diusion in the gas phase.27 maximum Tp and Tc values, which exceed room temperature, Deviation of the (Ln+M)/M ratio from the stoichiometric conform to M=Sr at the optimal doping level x=0.3.Our value in Ln1-xMxMnO3 results in coarsening of the surface data (Fig. 6) are in agreement with the prediction in ref 12. morphology even at a thickness d#1000 A° .La0.7Sr0.3MnO3/LaAlO3 shows a metal-like temperature dependence of the resistivity in the range 77–300 K with a Electrical properties sharper increase of the resistivity as T approaches 300 K. La0.7Ca0.3MnO3 reveals a resistivity maximum at 230 K. It was recently demonstrated for bulk samples4,10,12 that for La1-xMxMnO3 the P�FMM transition temperatures (both This value practically coincides with the data for the bulk material of the same composition.10 This value is also com- the resistivity maximum temperature, Tp, and the Curie temperature, Tc) can be presented in the form of a universal rA parable with Tp=220 K for a La0.72Ca0.25MnO3/MgO thin film prepared by magnetron sputtering,28 Tp=250 K for a dependence (certainly, at a constant Mn4+/Mn3+ ratio).The J.Mater. Chem., 1997, 7(5), 747–752 749Fig. 7 Temperature dependence of the resistance of La0.7Ca0.3MnO3 Fig. 6 Temperature dependence of the resistance of the films on films on LaAlO3 and SrTiO3 LaAlO3 (abbreviations are explained in Table 2) La0.67Ca0.33MnO3/LaAlO3 thin film prepared by laser normalization to maximum resistivity they practically match each other after a 20 K shift along the temperature axis.Taking ablation and annealed in oxygen at 1000 °C,29 Tp=270 K for La0.8Ca0.2MnO3/LaAlO3 and Tp=245–275 K for into account the lattice parameters of the substrates one can see that they possess nearly equal but opposite lattice mismatch La0.67Ca0.33MnO3/LaAlO3 films prepared by MOCVD.21,30 On the other hand, Tp and Tc values for La1-xCaxMnO3 with the film.Furthermore, according to ref. 21, Tp values of the epitaxial films on LaAlO3 and NdGaO3 [the latter has a films prepared by laser ablation are known to be strongly influenced by variation of post-deposition annealing: for very low lattice mismatch with the film (ca. 0.3%)] are very close. Thus, a tension but not a constriction of the film in the instance, Tp#95 K and Tp>300 K were reported for La0.67Ca0.33MnO3/LaAlO3 .19,29 One can suppose that vari- interface plane results in the change of Tp.An HREM study is desirable to understand which way the tension at the ation of the Mn4+/Mn3+ ratio occurs in such films owing to: (1) perhaps a broader d range between the reduced and interface influences the properties of the whole film. Along with variation of rA by changing the M cation, one oxygenated stas than that for La1-xSrxMnO3; (2) structural disorder induced in the film at the preparation stage; (3) devi- can also vary rA value by changing Ln.Fig. 6 shows that in this case the electrical properties do not vary gradually with ations of cation stoichiometry, especially the La+Ca/Mn ratio. The exact reason for this variation is not yet known.rA. In fact, the resistivity curves for La0.7Sr0.3MnO3 and La0.35Pr0.35Sr0.3MnO3 are similar, while Pr0.7Sr0.3MnO3 In ref. 19 it was suggested that epitaxial quality of the films is responsible for the shift of Tp. This suggestion seems rather reveals quite dierent behaviour. Certainly, Tp is expected to diminish with decreasing rA, but gradually.10,12 unlikely, taking into account the proximity of Tp values listed above for ceramics and thin film materials prepared by dierent As was mentioned in the introduction, the substitution of La for Pr or Nd leads to much more drastic changes in the techniques, and the dierent epitaxial quality.Next, anomalous Tp were observed only for films as-grown under low pO2 or electrical and magnetic properties than one would expect from the variation of rA.Study of bulk materials points to the after high-temperature (ca. 1000°C) annealing in oxygen. In both cases variation of the Mn4+/Mn3+ ratio is much more formation of charge- and spin-ordered phases. Let us try to construct a conceptual phase diagram of likely than a change of the epitaxial quality. It is worth mentioning that it is incorrect to assume that higher r means Ln1-xMxMnO3.Restricting the consideration by collinear spin ordering, we have ferromagnetic or antiferromagnetic states as better epitaxy.19 In fact, the metallic conductivity should be of the same order of magnitude for all compositions, but the well as a spin-disordered paramagnetic state. Superposition of the possible charge ordering (in other words, the ordered array resistivity of the P phase increases exponentially if this phase expands its stability field to lower temperature.Hence, of Mn4+ and Mn3+ ions31 ) and magnetic ordering increases the number of phase states by a factor of 2: 3×2=6 states. La1-xBaxMnO3 or La1-xSrxMnO3 films can be of the same epitaxial quality as La1-xCaxMnO3 ones, but they still have The relative positions of the phase fields in the coordinates temperature–magnetic field can be predicted in general: the higher Tp values.Another remarkable feature of La1-xCaxMnO3 films was paramagnetic phase occupies the whole high-temperature area, the antiferromagnetic state tends to low temperature and zero described in ref. 21. The Tp value was found to be dependent on the substrate material: an La0.8Ca0.2MnO3 film on MgO field, charge ordering is stabilized by decreasing temperature, and the ferromagnetic state is reinforced by the magnetic field. has Tp=240 K and a broader transition width than a film on LaAlO3 (Tp=270 K). The origin of the eect is not clearly Less evident but proven13,15 for Ln1-xMxMnO3 is the fact that a magnetic field destroys charge-ordered states.These understood yet: on the one hand, since the lattice mismatch between the film and the substrate has the opposite sign for relationships are illustrated by Fig. 8 which turns out to be a good guide for comparing the properties of Ln1-xMxMnO3. LaAlO3 and MgO and is significantly larger for MgO, strain can be induced in the film superimposing a deformation on Depending on Ln and M, the position of the zero field line (or the T -axis) changes.In high field only FMM and P states the three-dimensional array of the MnO6 octahedra, while on the other hand, since the film on MgO was not purely epitaxial are present. Thus, in high field, Ln1-xMxMnO3 with small rA would approximate the behaviour of La1-xSrxMnO3 in zero [a low-intensity (110) reflection was observed in the XRD diraction pattern21], the epitaxial quality could influence Tp.field. For instance, Pr0.65Ca0.35MnO3 requires H#12 T for such a transformation. Comparison of the La1-xCaxMnO3 In order to clarify this eect we compared La0.7Ca0.3MnO3 films on LaAlO3 and SrTiO3. According to the XRD charac- and Pr1-xCaxMnO3 positions in the conceptual phase diagram leads to the idea that the solid solution of the phases will be terization both films were of the same epitaxial quality. Resistivity curves are very similar for both films (Fig. 7); after subjected to such a transformation in lower field; in other 750 J. Mater. Chem., 1997, 7(5), 747–752Fig. 8 Conceptual phase diagram of Ln1-xMxMnO3: P, paramagnet; COI, charge-ordered state of paramagnet; FMM, ferromagnet with metal-like conductivity; CFMI, charge-ordered ferromagnet; AF, antiferromagnet; CAFI, charge-ordered antiferromagnet.Dashed line indi- Fig. 9 Temperature dependence of the resistance of cates the position of the T -axis (zero magnetic field) floating depending La0.35Pr0.35Ca0.3MnO3/LaAlO3 for dierent magnetic field values on Ln and M.removal of the field. In ref. 14 it was indirectly shown that words, increasing rA (up to that of La1-xSrxMnO3) and ordering of Pr moments may be responsible for this eect. increasing the field have equivalent influences. The La0.35Pr0.35Ca0.3MnO3/LaAlO3 specimen diers from It is noteworthy that dierent phase fields in the diagram Pr1-xCaxMnO3 in two important aspects: (1) the FMM phase dier drastically in resistivity and its temperature dependence. formed spontaneously under heating is actually stable in the High conductivity is possible only in the case of ferromagnetic range 20–70 K; (2) the hysteresis takes place without an ordering due to spin coupling of the itinerant (derived from eg external magnetic field.The first distinction agrees with the states) and localized (derived from t2g states) electrons of Mn conceptual phase diagram discussed above.The second ions. If the atomic spins of the Mn ions are disordered or distinction means that one more phase transition occurs at ordered antisymmetrically, then electronic exchange will be about 20 K in La0.35Pr0.35Ca0.3MnO3/LaAlO3 . By analogy suppressed.In the case of antiferromagnetic order the resistivity with Pr1-xCaxMnO3 one can suppose the ordering of Pr should be higher than in the case of paramagnetic disorder, as moments, but spontaneously. Certainly, some other mechanism between disordered atomic spins some nearly parallel spins may be possible. At any rate, the phase transition can explain occur, hence short-range exchange of electrons is possible. The why heating after cooling to 4.2 K results in a low resistivity antiparallel ordering of spins is accompanied by charge window at 20–70 K, but heating after cooling to 60 K does ordering in Ln1-xMxMnO3,13,15 in turn the charge-ordered not (at 60 K resistivity tended to immeasurably large values).phase derived from the paramagnetic state tends to result in In accordance with its expected position in the conceptual an antiferromagnetic-like temperature dependence of the mag- phase diagram, La0.35Pr0.35Ca0.3MnO3/LaAlO3 turns out to netization.32 Thus, the phase transitions from these ordered be very sensitive to the magnetic field.In a magnetic field of states to the FMM state give rise to a drastic change of the 1 T the cooling and heating resistivity curves are very close.resistivity. Being induced by a magnetic field, they produce r Only the FMM phase can be detected below 75 K in this field. values much larger than those caused by the influence of the As a result very large r values were found below 21 K as well magnetic field on the FMM�P transition. Therefore, phases as around 70 K. Increasing the field to 3 T results in a nearly in which the zero field line intersects CAFI, AF or COI phase saturatedGMReect (Fig. 10). The R–H hysteresiswas studied areas (Fig. 8), are expected to demonstrate very large r values at 75 K in more detail. Curves obtained by increasing and in a moderate field. Such a situation occurs at about half decreasing the field dier significantly, which is characteristic height of the conceptual phase diagram. for the COI�FMM transition15 and corresponds to crystalliz- This conclusion is well supported by experiments on ation and melting of charge-ordered state.Note that the Nd0.47Sm0.03Sr0.5MnO3,32 Pr0.7Ca0.25Sr0.05MnO3 11 and very recently La1/3Nd1/3Ca1/3MnO3.33 Here we report a new example. A thin film of La0.35Pr0.35Ca0.3MnO3 on LaAlO3 reveals a very complicated temperature dependence of the resistivity (Fig. 9). The magnetic measurements were not accomplished in the work; consequently the assignment of the magnetic states can be only preliminary and based on analogy. 13–15 The behaviour of the film resistivity was dierent depending on which process (cooling or heating) was used. Upon cooling the resistivity increased exponentially and became immeasurable below 70 K.After cooling to 4.2 K the specimen was heated, and it demonstrated a sharp drop of resistivity in the range 20–70 K (Fig. 9). Above 80 K the cooling and heating curves converged. The hysteresis eects observed earlier13–15 were explained by charge ordering of the P phase (probably overlapped by canted aniferromagnetism).As the result, the first-order phase transition FMM�COI takes place. For Pr1-xCaxMnO3 (x#0.3–0.4) only the COI phase is stable in zero field. In a magnetic field of ca. 4–5 T at temperatures<100 K the firstorder phase transition was induced. At temperatures below Fig. 10 Resistance, R, vs. magnetic field, H, loops for La0.35Pr0.35Ca0.3MnO3/LaAlO3 : the path direction is a-b-c-d 25–30 K the metastable FMM state could be frozen after J.Mater. Chem., 1997, 7(5), 747–752 751measurement data in the literature points to an interdependence of the phenomena. (c) Non-collinear antiferromagnetism can be suspected in many particular cases. At any rate, the idea that compositions most interesting for practical applications (owing to low saturation field, high r and dr/dH) should be searched at about the half-height level of the conceptual phase diagram, provides a good guide for further studies.This research was partly supported by Volkswagen Foundation project I/69341 and RFBR grant no. 96–03–33027. References 1 C. N. R. Rao and A. K. Cheetham, Science, 1996, 272, 369. 2 M. H. Kryder, W. Messner and L. R. Cartley, J. Appl.Phys., 1996, 79, 4485. 3 T. Yotsuya, Jpn. J. Appl. Phys., 1996, 35, L23. Fig. 11 Temperature dependence of the resistance of 4 R. Mahesh, R. Mahendiran, A. K. Raychaudhuri and C. N. R. Rao, La0.35Pr0.35Ca0.3MnO3/LaAlO3 cooled at H=1 T to 4.2 K with J. Solid State Chem., 1995, 120, 204. subsequent heating to 160K and cooling in zero field 5 R. von Helmolt, J. Wecker, B. Holzapfel, L.Shultz and K. Samwer, Phys. Rev. L ett., 1993, 71, 2331. 6 H. L. Ju, J. Gopalakrishnan, J. L. Peng, Q. Li, G. C. Xiong, ‘crystallization’ occurs in a very low field, ca. 0.2 T; nevertheless T. Venkatesan and R. L. Greene, Phys. Rev. B, 1995, 51, 6143. 7 R. Mahendiran, A. K. Raychaudhuri, A. Chainani, D. D. Sarma the resistivity in zero field was reproduced after cycling in a and S.B. Roy, Appl. Phys. L ett., 1995, 66, 233. field of the opposite orientation. One can assume that careful 8 A. Urushibara, Y. Moritomo, T. Arima, A. Asamitsu, G. Kido and variation of the temperature at about 70 K can provide an Y. Tokura, Phys. Rev. B, 1995, 51, 14103. even lower saturation field as well as a thinner hysteresis loop, 9 P. Schier, A. P. Ramirez, W. Bao and S.W. Cheong, Phys. Rev. owing to the fact that the FMM phase is stable below 70 K. L ett., 1995, 75, 3336. 10 H. Y. Hwang, S. W. Cheong, P. G. Radaelli, M. Marezio and It should be mentioned that the reproduced resistivity value B. Batlogg, Phys. Rev. L ett., 1995, 75, 914. in zero field after cycling in the magnetic field was several 11 B. Raveau, A. Maignan and V. Caignaert, J.Solid State Chem., times lower than the initial zero field resistivity. This result 1995, 117, 424. may imply some remanent magnetization in the sample. An 12 P. G. Radaelli, M. Marezio, H. Y. Hwang and S. W. Cheong, irreversible change of the magnetic state was assumed to be J. Solid State Chem., 1996, 122, 444. 13 Y. Tomioka, A. Asamitsu, Y. Moritomo, H. Kuwahara and an alternative explanation in ref 14, where similar behaviour Y.Tokura, Phys. Rev. B, 1996, 53, 1689. had been registered for Pr0.6Ca0.4MnO3, but no remanent 14 M. R. Lees, J. Barratt, G. Balakrishnan, D. McK. Paul and magnetization was detected by magnetic measurements. M. Yethiraj, Phys. Rev. B, 1995, 52, 1. Similar to Pr1-xCaxMnO3, the FMM state in 15 H. Kuwahara, Y. Tomioka, A. Asamitsu, Y.Moritomo and La0.35Pr0.35Ca0.3MnO3/LaAlO3 can be frozen in zero field Y. Tokura, Science, 1995, 270, 961. 16 Y. Tomioka, A. Asamitsu, Y. Moritomo, H. Kuwahara and after cooling in the magnetic field (even for H=1 T). Y. Tokura, Phys. Rev. L ett., 1995, 74, 5108. Subsequent heating to 160 K in zero field produces a resistivity 17 G. H. Jonker, Physica, 1956, 22, 707. curve which looks like that one produced in a field (Fig. 11). 18 G. C. Xiong, Q. Li, H. L. Ju, S. N. Mao, L. Senapati, X. X. Xi, The next cycle of cooling from 160 K leads to resistivity R. L. Greene and T. Venkatesan, Appl. Phys. L ett., 1995, 66, 1427. behaviour as described above for the sample which was not 19 S. Jin, T. H. Tiefel, M. McCormack, R. A. Fastnacht, R. Ramesh and L. H. Chen, Science, 1994, 264, 413.exposed to the magnetic field, except for in the low-temperature 20 K. Li, L. Liu, J. Sun, X. J. Xu, J. Fang, X. W. Cao, J. S. Zhu and range. Below 35 K the resistivity drops again to a value about Y. H. Zhang, J. Phys. D, 1996, 29, 14. 103× higher than the resistivity of the frozen FMM phase 21 Y. Q. Li, J. Zhang, S. Pombrik, S. DiMascio, W. Stevens, Y.F. Yan and was nearly constant below 20 K. This trend occurs for and N. P. Ong, J.Mater. Res., 1995, 10, 2166. the subsequent heating too, but above 25 K the resistivity 22 F. Weiss, K. Fro� hlich, R. Haase, M. Labeau, D. Selbmann, J. P. Senateur and O. Thomas, J. Phys., Colloq., 1993, 3, 321. again approximates that of the sample not exposed to the 23 O. Yu. Gorbenko, V. N. Fuflyigin, Y.Y. Erokhin, I. E. Graboy, field. It is noteworthy that Pr0.6Ca0.4MnO3 reveals similar A. R. Kaul, Yu. D. Tretyakov, G. Wahl and L. Klippe, J. Mater. behaviour in a field of 4 T.14 We see again, as for the sample Chem., 1994, 4, 1585. not exposed to the magnetic field, that low-temperature 24 J. A. M. van Roosmalen, P. van Vlaanderen, E. P. H. Cordfunke, reordering (probably with a contribution from Pr W.L. Ijdo and D. J. W. Ijdo, J. Solid State Chem., 1995, 114, 516. 25 J. A. M. van Roosmalen and E. P. H.Cordfunke, J. Solid State moments), which occurs in Pr1-xCaxMnO3 in the field, for Chem., 1994, 110, 109. La0.35Pr0.35Ca0.3MnO3/LaAlO3 is spontaneous. 26 I. G. Krogh Andersen, E. Krogh Andersen, P. Norby and E. Skou, J. Solid State Chem., 1994, 113, 320. 27 C. H. J. Van den Brekel and A. K. Jansen, J. Crystal Growth, 1978, Conclusions 43, 364. 28 K. Chahara, T. Ohno, M. Kasai and Y. Kozono, Appl. Phys. L ett., The results obtained by us and described in the literature for 1993, 63, 1990. Ln1-xMxMnO3 are in good agreement with the proposed 29 M. Jaime, M. B. Salamon, K. Pettit, M. Rubinstein, R. E. Treece, conceptual phase diagram. There is an analogy between the J. S. Horwitz and D. B. Chrisey, Appl. Phys. L ett., 1996, 68, 1576. magnetic field and rA increase eects. Nevertheless the follow- 30 G. J. Snyder, R. Hiskes, S. DiCarolis, M. R. Beasley and ing specific points are worthy of note. T. H. Geballe, Phys. Rev. B, 1996, 53, 14434. 31 J. Goodenough, Phys. Rev., 1955, 100, 564. (a) If the ionic radii of cations in the A position vary 32 H. Kuwahara, Y. Tomioka, Y. Moritomo, A. Asamitsu, M. Kasai, significantly, then local stress may be important, which is the R. Kumai and Y. Tokura, Science, 1996, 272, 80. case for Ba-containing Ln1-xMxMnO3. Next, the magnetic 33 G. H. Rao, J. R. Sun, J. K. Liang, W. Y. Zhou and X. R. Cheng, moment of rare-earth-metal cation cannot be neglected at low Appl. Phys. L ett., 1996, 69, 424. temperature. (b) Antiferromagnetism and charge ordering were considered Paper 6/06465E; Received 19th September, 1996 mainly independently in this study. Nevertheless the magnetic 752 J. Mater. Chem., 1997, 7(5), 747&ndash

ISSN:0959-9428    

DOI:10.1039/a606465e

出版商:RSC    

年代:1997

数据来源: RSC