Materials with ultrafine microstructures: Retrospectives and perspectives

NanoSTRUCTUREDMATERIALS VOL. 1, PP. 1-19, 1992 COPYRIGHT©1992 PERGAMONPRESSplc ALL RIGHTSRESERVED

0965-9773/92 $5.00 + .00 PRINTED IN THE USA

MATERIALS WITH ULTRAFINE MICROSTRUCI'URES: RETROSPECTIVES AND PERSPECTIVES H. Gleiter, Universi~t des Saarlandes und Institute fiir Neue Materialien, Geb~iude 43, W-6600 Saarbriicken, Germany

Introduction Materials with ultrafine microstructures are solids that contain such a high density of defects (point defects, dislocations, grain boundaries, interphase boundaries etc.) that the spacings between neighboring defects approach interatomic distances. Ultrafine microstructures have already been recognized more than a century ago (e.g. by Sorby, Tschernoff, Osmond, Howe, Sauveur, Wilm, Merica et al.) to exhibit remarkable and technologically attractive properties. The physical reasons for these properties were elucidated when modern defect theory and high resolution methods (TEM, FIM) became available. Since about 1970, the interest of materials scientists refocused on solids with ultrafine microstructures when it was recognized that specifically tailored ultra.fine microstructures permit the generation of materials with new atomic and/or electronic structures. So far, the following two developments of this type have emerged. In the area of semiconductors, quantum well structures and supedattices consisting of thin layers with different dopings or compositions were generated by means of two-dimensional ultrafine microstructures. Research activities aiming towards the synthesis of new atomic (and electronic) structures in metals, ceramics and semiconductors by means of threedimensional ultrafine microstructure were initiated by the proposal to generate solids, a large volume fraction of which consists of the cores of grain and/or interphase boundaries. These solids have been called nanocrystalline or nanophase Solids. Such solids differ structurally from crystals and glasses with the same chemical composition, because the atomic arrangements formed in the cores of grain or interphase boundaries deviate from crystalline or glassy structures. Recent studies of nanocrystalline solids by x-ray/neutron diffraction, EXAFS, different spectroscopies, as well as property measurements support these ideas. Presently, the following areas of research on nanocrystalline materials appear to be of particular interest: nanocrystalline alloys, the ductility of nanocrystalline ceramics, ferromagnetism in nanocrystalline metals, nanoglasses and submicron materials. In this paper, a material will be called a material with an ultrafine microstructure if it is a solid and if it contains such a high density of defects (point defects, dislocations, grain boundaries, interphase boundaries, etc.) that the spacings between neighboring defects approach interatomic distances. Macromolecular materials and solids the atomic structure and/or properties of which are dominated by free surface effects (e.g. thin films, thin wires, small isolated crystals) will not be considered. According to the title, the paper will be divided in the following two sections: (i) retrospective and historical development, and (ii) recent developments and future perspectives. Retrosoective and Historical Develovment Historically, two periods of development in research and application of ultrafine microstructures may be distinguished. During the first period, ranging from about 1870 to about 1970, the microstructure of materials

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was recognized to be the crucial parameter controlling many mechanical, magnetic and electronic properties. This line of thought seems to start with the pioneering work on the mechanical properties of iron alloys by Sorby, Tschernoff, Osmond, Howe, Sauveur and others before the turn of the century. Their studies led to the conclusion that the fine-scale microstrncture retained after the allotropic transformation of iron alloys gives martensite its hardness. The discovery of precipitation hardening by Wilm in 1906 was the first observation suggesting that the correlation between microstructure and properties (originally proposed for ferrous alloys only) applies to non-ferrous materials as well. Wilm quenched an AI-Cu-Mg-Mn alloy and noticed (after a long weekend) a substantial hardening relative to the as-quenched state. However, it was not until 1919 when Merica, Waltenberg and Scott proposed that the hardening resulted from precipitation of a new phase on a submicroscopic level. Numerous observations in the subsequent years substantiated and generalized this view and led to the classification of the properties of solids with different types of chemical bonding into microstructure-sensitive and non-sensitive ones (1). In the following years, this classification played an important role in promoting the idea of lattice defects and their significance for crystalline solids. In fact, the physical understanding of the mechanism by which ultrafine microstructures affect the properties of solids received a remarkable boost after World War II, from the advent of the theory of lattice defects--in particular, dislocation theory-and from the availabilityof new high resolution research techniques such as electron and field ion microscopy. Both developments helped to elucidate the physical basis for understanding the correlation between the structure sensitive properties and the micrnstructure of solids. As a matter of fact, the development of most high strength and high temperature materials available today is based on the results of those studies. When it was recognized that dislocations play a similar role for mechanical behavior of materials, as do domain wails or flux lines for ferromagnetic or superconducting properties, respectively, it became apparent that ferromagnetic and superconducting properties can be manipulated, too, by suitably varying the microstructure (2). In fact, the pinning of ferromagnetic domain walls or of flux lines in type II superconductors by finely dispersed precipitates leads to magnetic materials with high coercive forces and superconductors with high critical current densities (2). Instead of reducing the motion of dislocations by precipitates from supersaturated solid solutions, small second phase particles introduced by means of powder metallurgy or by extrusion of two phase coarse grained structures (3), (4) may be used as well. The enhanced defect density in irradiated, in highly cold worked (5), as well as in fine-grained materials (6) results in similar effects, since the defects inhibit the dislocation motion. A closely related approach in the field of ceramics generated by the sol-gel method was the proposal to use heterogeneity on a nanometer scale (7) (12). This approach was based on the hypothesis that diphasic or multiphasic ("nanocomposite') xerogels which are heterogeneous on a nanometer scale store more energy than a single phase gel and thus exhibit new properties. About 1970 the second period of developments in the area of ultrafine microstructures started, when it was recognized that specificallytailored ultrafine microstructures permit the generation of solids with new atomic and/or electronic structures. These developments seem to bring Feymnan's dream of nanotechnology (13) closer to reality. Up to now, the following two types of solids with new electronic and/or atomic structures have been generated by means of tailored ultrafme microstructures. In the area of semiconducting materials, extensive research activities were initiated on a new two-dimensionai class of materials with a nanometer-scale microstrncture by the proposal of an "engineered" semiconductor superlattice (14), (15) consisting of alternate thin layers with different dopings, different composition (Fig. 1) leading to new electronic structures. These supeflattices represent an extension of double- or multibarrier structures where quantum effects prevail because the layers have a thickness of a few lattice constants. Several recent reviews of the present state of understanding of the physics and application of quantum well structures and superlattices have been compiled in ref. (16). A recent modification and extension of this approach is the alternating growth of coherent multilayer structures of components with large lattice parameter mismatch. This opens the possibility of combining superlattice effects with those associated with strain to design matefiais of suitable band gaps and electronic properties. By growing thin strained layers of high crystalline quality, one obtains binary materials, the so-called pseudoailoys. PseudoaUoys permit to combine strain effects with quantum size effects. Twodimensional semiconductor materials systems exhibit the technolngically important features of high quantum efficiency and rapid carrier capture efficiency. An example of such a structure is a thin pseudomorphic layer of InAs embedded between In0.sa.M0.47Aslayers which are lattice-matched to an InP substrate (17)-(19). These

ULTRAFINEMICROSTRUCTURES

3

DOPING SUPERLATTICE CONDUCTION BAND

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VALENCE BAND COMPOSITIONAL SUPERLATTICE CONDUCTION BAND

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.J

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DISTANCE x

Fig. 1

Spacial variations of the conduction and valence band edges in two types of superlattices: doping (top) and compositional. Esl and Es2 indicate the gap energy in the different layers of the compositional superlattices. 1 is the superlattice periodicity.

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Fig. 2

1.0

22 Inis

W

Ino.53Gi

1.2

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1.6

2.95

Wavelength (pm)

Photoluminescent spectra at 4 K of an InAs quantum well 1 run and an InAs quantum well 3 nm thick embedded in lno.53Alo.,TAs barriers. The peaks labeled B and W correspond to emission from the barrier and quantum well, respectively. The two arrows indicate the values of the band-gap energies of the unstrained bulk In0.53Gao.,TAs and InAs (17).

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Fig. 3

Edge dislocation in a simple cubic crystal. In the dislocation core, the atomic density and coordination is changed in comparson to the perfect cubic lattice. In fact, the core of the dislocation may be interpreted as a channel formed by seven atoms as indicated in Fig. 3.

Fig. 4a

Atomic structure in the core of grain boundaries between two crystals tilted relative to one another by 36.9° about a common (100) direction normal to the plane of the figure. The structure was deduced (23) from the high resolution electron micrograph shown in Fig. 4b (section A). The boundary core structure may be described as two-dimensional periodic array of two different pentagonal polyhedra (indicated on the right side of the boundary by broken lines).

ULTRAFINEMICROSTRUCTURES

5

heterostructures contain a thin strained InAs quantum well. The strain modifies the properties of the resulting material. For example, the wave-length of the light emitted by the InAs layer is shifted from its usual bulk value of 2.95 ~tm to the range of 1.2-1.6/~m depending on the layer thickness (Fig. 2). Optical investigations of highly strained (3.4% lattice mismatch) isolated InAs quantum wells show not only a controllable emission in the 1.21.6ttm range but also an increase in the photoluminescence, by a factor of 2.5, relative to the bulk InAs value. Research activities aiming toward the generation of new atomic and electronic structures by means of ultrafine microstructures were stimulated by a proposal (20) to generate solids (metals, semiconductors, ceramics) a large volume fraction of which consists of the cores of grain (or interphase) boundaries. Solids of this type were argued to differ structurally and propertywise from the crystalline or glassy state with the same chemical composition because grain boundaries and interphase boundaries represent a special state of solid matter due to the fact that the atoms in the core of an interface are known to be subjected to the (periodic) potential field of the crystals on both sides of the interface (c.f. Figs. 4a and b). As a consequence, the arrangements formed by the atoms in the cores of interfaces differ from the glassy and the crystalline state. In the glassy state, the atoms are not subjected to the (periodic) boundary conditions like the atoms in the core of interfaces. In the crystalline state, the atoms relax to a structure of lowest free energy which is prevented in the cores of interfaces by the fact that the interface represents the transition between two crystals of different orientations. Hence, the atomic structure and the properties of solids in which the volume fraction of the cores of interfaces becomes comparable to or larger than the volume fractions of the residual crystals differ from the structure and properties of the same materials in the crystalline or in the glassy state. Direct experimental and theoretical evidence supporting this idea comes from studies of the structure and properties of defect cores, in particular, of grain and/or interphase boundaries (21), (22). Figs. 3 and 4 illustrate the atomic arrangements in the vicinity of an edge dislocation and a grain boundary between two cubic crystals. The cores of both defects are characterized by a reduced atomic density and a modified atomic structure in comparison to the perfect lattice. In fact, the core of the edge dislocation may be described in terms of a channel comprising seven atoms in every cubic plane (Fig. 3). The structure of the grain boundary core, shown in Fig. 4a, represents a periodic sequence of two different pentagonal polyhedral units. The physical reason for the reduced atomic density and the modified atomic arrangement may be seen by considering, for example, the structure of the grain boundary core shown in Fig. 4a. The boundary core is the region where two crystals with different crystallographic orientations are joined together. Due to the different orientations of the crystals, the two crystal lattices match poorly in the core region. The poor matching reduces the atomic density in the core regions by typically 10 to 30% relative to the perfect lattice. In other words, the density in the cores is considerably lower than in glasses. Furthermore, as the boundary core is the region of contact between two differently oriented crystals, its atomic structure depends on the orientation relationship between both crystals and the inclination of the boundary plane (21) (22). In other words, each time the orientation relationship and/or the inclination of the boundary plane is varied, a new boundary core structure results.1 Hence, in the various boundaries of a polycrystal (e.g. of Cu) consisting of many randomly oriented crystals, a large variety of boundary structures with different interatomic spacings are formed. However, the large variety of boundary core structures and the reduced densities are commonly not noticed in coarse grained polycrystals (crystal size >_ 10/~m) because the volume fraction of the boundary cores is 10-2% or less. Nevertheless, precise measurements (20), (24)-(26) do reveal the existence of both boundary core properties. On the other hand, if the volume fraction of the boundary cores is enhanced to about 50% or more by reducing the crystal size to a few lattice constants (i.e. a few nanometers), one obtains materials2 (Fig. 5) which consist structurally of the following two components of comparable volume fractions. tAn example may be seen in Fig. 4b. The boundary plane in section B is displaced relative to A. This displacement results in different boundary core structures in A and B. aPolycrystals with an average crystal size of a few nanometers will be called in the subsequent section nanocrystalline materials.

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Fig. 4b

Electron micrograph of a symmetric 36.9° (100) tilt grain boundary in NiO (23). The black regions represent the atomic positions. In the boundary core the atomic density and coordination is changed in comparison to the perfect lattice. The boundary structures in the facets A and B are different due to the vertical displacement of the boundary plane relative to the lattices of the two crystals.

Fig. 5

Atomic structure of a two-dimensional nanocrystalline material. The structure was computed by the following the procedure given in ref. 27. The interaction between the atoms was represented by a Morse potential fitted to gold. The atoms in the centers of the "crystals" are indicated in black. The ones in the boundary core regions are represented by open circles.

7

U LTRAFINE MICROSTRUCTURES

50

I

25-

0 x -25 ~

-50 I

I

5

10

15

k IA'~I

1500 --

~.,~1000 -

500

00

2

/. 6 8 INTERATONIC DISTANCES r [~,]

10

Fig. 6a,b The weighted EXAFS oscillations Xk a (Fig. 6a) and the corresponding Fourier transform FT (Xka), (Fig. 6b, phase shift is not included) of a nanocrystalline Cu sample ( ) (crystal diameter: 10 nm) in comparison to polycrystalline Cu (+ + +). In the nanocrystalline sample the amplitude of the EXAFS oscillations and the FT (Xk 3) are reduced relative to the polycrystaUine Cu (34). The numerous, small crystals with different crystallographic orientations forming the "crystalline component" and a network of boundaries the structure of which differs from boundary to boundary. According to the boundary properties discussed above, the structural component formed by all boundaries (called the boundary component) has the following two features:

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(i) The average atomic density in the boundary regions is reduced by 10 to 30% relative to the crystal density depending on the type of chemical bonding between the atoms. The boundary density cannot be enhanced by inserting additional atoms into the boundary because the local boundary free volume (i.e. the local free volume dispersed between the boundary atoms) is smaller than the atomic volume so that interatomic penetration would result in the region where an atom would have been inserted. For example, the free volume in the center of the two pentagonal polyhedra (Fig. 4a) is less than one atomic volume. (ii) The boundaries exhibit a broad spectrum of interatomic spacings ranging from closely packed atoms to widely spaced ones. The wide spectrum of interatomic spacings cannot be removed by atomic relaxations. This may be seen by considering the interatomic spacings in the boundary core of Fig. 4a in comparison to the spacings in the crystal lattices.

I

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~.~ .?-..',.',- -.~;,,-,,~

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J 0.80

0.60

I -6

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VELOCITY

Fig. 7

I 3 [mm/s]

I

0.5

0S

[mm/=]

The coarse-grained M6ssbauer spectra and quadrupole splitting (Q.S.) distributions for (a) FeF2 powder and (b) nanocrystalline FeF2 (77).

The reduced density in the boundary cores, as well as the wide distribution of interatomic spacings seems to be born out by experimental observations. High resolution microscopy (23), and grazing incidence diffraction of x-rays on Au-bicrystals (24) (28) indicate a reduction of the atomic density in the core region between 15 and 40%. Similar numbers are suggested by recent theoretical studies on boundaries in bicrystals (76). Small angle scattering measurements on nanoerystalline Pd (20), (30), and TiOa (31) indicated the average density in the boundary cores (width between 0.5 and 1 nm) to be reduced by about 20 to 40%. Information about the atomic

ULTRAFINEMICROSTRUCTURES

9

arrangement in the boundary cores of nanocrystalline materials was obtained by static and dynamic studies. EXAFS measurements (34) indicate a reduction of the number of nearest, next nearest, etc. neighbors averaged over all atoms of a nanocrystalline material relative to a single crystal (Figs. 6a and 6b). This result may be understood, if the structure in the vicinity of the atoms in the boundary cores differ from the structure in the lattice. In fact, the reduction observed suggests a fraction of the boundary atoms to be displaced from the sites they would occupy in a perfect lattice. A similar conclusion is suggested by Mtssbauer spectroscopy measurements on nanocrystalline FeFz (69). The quadrupole splitting distribution of nanocrystalline FeFz is broadened (Fig. 7) relative to the one of an FeFz single crystal indicating a variety of atomic configurations with non-lattice symmetry to exist in the boundary cores. X-ray diffraction studies (33) on nanocrystalline Fe have been interpreted in terms of a boundary core structure in which the atoms are displaced from the ideal lattice sites of both crystals (Fig. 8).

i0

v

0.2 0.4 0.6 0.8 1.0 1.2 1.4 1.6 wave

F~g. 8

vector

2sin@/l {A")

Comparison of the measured (-- + --) and computed ( ) interference functions of nanocrystaUine Fe. The model computed is a mixture of 6-rim (75 vo1.%) and 4-nm (25 vol.%) Fe crystals in which the boundary atoms are displaced from the ideal lattice sites in random directions. The displacement distances were chosen as follows. The atoms in the boundary core were displaced by 50% and the two subsequent layers by 25% of the nearest neighbor distance relative to the lattice sites they occupied originally. The layers of displaced atoms were used to model the ensemble average of the structure of the boundary core. The atomic density in the core region was adjusted to the measured boundary density (30), (33).

These displacements result in a diffuse x-ray diffraction background. However, the absence of this background in the diffraction curves of nanocrystalline Pd or Cu suggests that the atomic structure in the boundary cores of nanocrystalline materials may depend on the chemical composition. This conclusion agrees with the well-established different atomic relaxations at free surfaces and in the dislocations cores of chemically different materials. In other words, in a nanocrystalline material the boundary core structures (and hence the properties of this material) may be manipulated by varying the chemical composition in the boundary cores, e.g. by incorporating solute atoms into the boundary cores. The modification of the boundary properties due to boundary segregation of solute atoms is well known for numerous alloys (e.g. the embrittlement of Cu by Bi segregation to the boundaries or the enhanced ductility of Ni3AI if B is present in the boundaries). The structural deviation between the crystals and the boundary cores are also born out by methods probing the dynamic displacement of the boundary atoms. Measurements of the Debye temperature by Mtssbauer spectroscopy indicate a reduction by about 100K in nanocrystalline Fe (77) and about 210K in nanocrystalline

10

H GLEITER

FeFz (78). The Debye temperatures of Fe and FeF2 are 467K and 298K, respectively. These results seem consistent with inelastic neutron scattering measurements (79) suggesting additional low frequency phonons in nanocrystalline Ni in comparison to a coarse grained polycrystal of the same material. The low frequency phonon modes may be the reason for the enhanced specific heat, %, (80), (81) of nanocrystalline materials (e.g. for 6nm Pd, % is enhanced relative to a coarse grained polycrystal by about 40% at 300K). Measurements of the elastic constants by bending, sound wave propagation or Brillouin scattering indicate a reduction of the Youngs and the shear modulus by 15% or more (82). The reduction cannot be caused by residual porosity because the results of the last two methods depend little or not at all on porosity. However, all of the observations reported may be understood by assuming (i) the average atomic density in the boundary cores to be reduced relative to the perfect crystal and (ii) this reduction to results in atoms that are more weakly coupled to their neighbors than the ones in the perfect lattice. Although the above discussion has been limited on manipulating the atomic structures of solids by means of incorporating a high density of grain boundaries, basically the same arguments hold for other lattice defects as well, such as dislocations, interphase boundaries, etc.. In recent years, the research activities on the production, the investigation of atomic structure and properties of nanocrystalline materials increased remarkably. The progress made has been documented in several review articles (32), (35), - (40). Long before materials with ultrafine microstructures have been synthesized by man, they have been formed in several natural composites. For example, human teeth are an exquisitely tailored composite at the 1-2 run level of fibrils of hydroxyapatite topotacticaUy grown onto collagen. A similar structure was found in natural corals. Semicrystalline polymers are composed of plate shaped crystalline regions with a thickness of typically 1-3 nm separated by "amorphous" regions consisting of tie molecules between adjacent crystallites, as well as macromolecules folded in a regular or irregular fashion. Probably the oldest solids with an ultrafine microstructure are primitive meteorites, in which aggregates of carbon atoms or silicon carbide of 5 nm or less are found. These aggregates are believed to be formed by condensations in early age of the solar system of interstellar space (41), (42). Asbestos, opals and calcedon represent minerals with an ultrafine microstructure formed on the surface of the earth many million years before any man-made substance of this type was available. Early applications of materials with ultrafine mierostructure have been revealed in Egypt and China, where the size of the pigments of certain colors has been reduced into the submicron regime by grinding or milling, in order to enhance the intensity of the colors. present Situation Up to now, the research activities on materials with ultrafine microstructures have resulted in several hundred publications on quantum well structures and superlattices. More than one hundred papers have been published dealing with the various aspects of the structure and the properties of materials containing a high density of defects, preferably grain boundaries. In order to keep this manuscript within reasonable length, we refer to existing review articles concerning the present state of the art (15), (32), (35) - (40). We shall now turn to the second part of the paper dealing with recent developments and future perspectives. Recent Developments and Persoectives In this section we shall point out some recent developments which may become the nuclei for new areas in the future. Naturally, the selection is bound to be subjective. Therefore, the author apologizes to all colleagues whose work will not be discussed. Alloys with Ultrafine Microstructures Alloys made up of nanometer-sized crystals with different chemical compositions have recently been synthesized (43) and seem to open the way to generate alloys of chemically different components that are insoluble in coarse-grained polycrystals. The preparation of nanometer-sized alloys was performed by means of an UHV evaporator filled with pure helium (~0.1 MPa pressure). Inside of the evaporator, several evaporation (or sputtering) units were arranged for the substances A, B, C... which were evaporated (or

11

ULTRAFINE MICROSTRUCTURES

sputtered) simultaneously. The small crystals of A, B, C ... formed in the He atmosphere were accumulated on the rotating cold finger as a loose powder. Under suitable evaporation (sputtering) conditions, the powder is a random mixture of small crystals of A, B, C .... Subsequent consolidation of the stripped-off powder lead to a nanometer-sized polycrystal, the individual crystallites of which consisted of A, B, C .... This product was called a "nanoerystalline alloy." Compared to single component nanoerystalline solids, nanocrystalline alloys contain a high density of interphase boundaries as well as grain boundaries ~. The atomic structure of nanoerystalline Ag-Fe alloys was studied recently (43). Ag-Fe alloys were selected for two reasons. Ag and Fe are insoluble in the solid and liquid state. Furthermore, the atomic structure may be investigated conveniently by x-ray diffraction as well as M6ssbauer spectroscopy. Fig. 9 displays the M6ssbauer spectrum of a nanocrystalline Ag-30 at% Fe alloy with an average crystal size of 8 nm. The spectrum consists of three components of comparable volume fractions. t . O0

C 0

0.99 E C L I--

0.98-

0 C_

÷

-12

I

I

I

I

I

-8

-4

0

4

8

velocity

Fig. 9

[mm/s]

MSssbauer spectrum of nanocrystalline Fe-Ag alloy (30at% Fe, 10 K, average crystal size 8 nrn). The spectrum consists of the following three components. (i) ~,-Fe ( ), (ii) Fe atoms dissolved in Ag (-.-) and (iii) Ag atoms dissolved in Fe (---) (43).

One component corresponds to body centered cubic (bcc) a-Fe crystals, the two other components originate from bee ,,-Fe crystals with incorporated Ag atoms, and face centered cubic (fee) Ag crystals with incorporated Fe atoms. In addition, the x-ray scattering data indicate the presence (Fe-free) Ag crystals. The atomic structure of the nanocrystalline Ag-30at%Fe alloy deduced from these observations is shown in Fig. 10. Apart from pure bcc ,~-Fe and pure fee Ag crystals, solid solutions are formed in the strained lattice regions in the vicinity of the Ag/Fe interphase boundaries and/or at the Fe/Fe, Ag/Ag, grain boundaries, respectively. An enhancement of the solute solubility due to elastic strains has been demonstrated in the past theoretically (44), (45) and experimentally (46) for various alloys. The alternative interpretation of the data reported above in terms of Fe or Ag atoms which become trapped in Ag or Fe crystals during the evaporation and crystallization process was ruled out by separating the Ag and Fe vapor sources. The formation of Ag/Fe 3Preparation of nanocrystalline materials and alloys under UHV conditions seems crucial to guaramee the formation of uncontaminated interfaces. In fact, nanometer-sized materials synthesized by using crystals with contaminated surfaces usually do not exhibit the features discussed in this paper. Of course, deliberate adsorption of specific atoms at the surface of the small crystals prior to compaction may be attractive because it permits to incorporate certain atoms (molecules) into the boundaries. As a matter of fact, this method has been used to dope the boundaries of nanocrystalline materials, e.g., doping of the grain boundaries in nanocryStalline F¢ or Cu by Bi.

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H GLEITER

solid solutions in Ag-Fe nanocrystaliine alloys seems remarkable because Fe and Ag are immiscible in the solid and liquid state. If solid solutions are formed in such an extremely insoluble system, it seems conceivable that nanocrystalline alloys may open the way to the formation of alloys in systems in which alloying has not yet been possible. Alloying of Ag-Fe by co-evaporation on cold substrates is possible. However, it is limited to thin films.

D~O0'u 0 O t J ~ ~ ~ ~" ) 0 0 0 0 0 D.. ~. ,~0 ~0 ~'~'~/~-~W2r~m~'~k-~C-~'-P~"~-~,. ,_" .c D

0 o

eccr( ...-

Fig. 10

Schematic model of a nanocrystalline Ag-Fe alloy according to the data of M6ssbauer spectroscopy (Fig. 9). The alloy consists of a mixture of nanometer sized Ag and Fe crystals. In the (strained) interracial region between Ag and Fe crystals solid solutions of Fe-atoms in Ag and Ag-atoms in Fe are formed, although both components are insoluble in the liquid and/or solid state. Similar effects may occur in the grain boundaries between adjacent Fe and Ag crystals (43).

Nanoglasses So far a frequently used procedure to synthesize nanocrystaUine solids is by consolidation of small crystals as was described in the previous section. By analogy, non-crystalline solids may be generated by consolidating nanometer-sized glassy spheres (47) - (49). The resulting solids were called "nanogiasses." The preparation follows basically the same route as was discussed for nanocrystalline alloys. However, instead of evaporating a material that crystallizes upon cooling in the He-atmosphere, one now evaporates (or sputters) a material that solidifies in the form of small glassy spheres. These spheres are subsequently consolidated into a nanogiass. The atomic structure proposed (as a working hypothesis) for nanoglasses in the as-consolidated state is indicated schematically in Fig. 11. The regions of contact between adjacent droplets differ structurally and/or chemically from the atomic structure in the center of the droplets for the following reason. The atoms at the surface of an isolated glassy droplet in vacuum form a glassy (short range ordered) atomic arrangement with the atoms in the interior so that certain interatomic spacings are preferred between the surface atoms and the atoms in the interior. If two (originally isolated) glassy droplets are brought into contact (e.g., in the region of contact between A and B in Fig. 11), the interatomic spacings between the surface atoms belonging originally to the droplets A and B deviate from the interatomic spacings between the atoms in the interior of A and B. Although some relaxational motion of the atoms in the region of contact may occur, density and interatomic spacings in the interfacial regions is likely to differ from a bulk glass with the same chemical composition. In fact, by analogy to the structure of nanocrystalline materials, the atomic structure of the interfacial regions between the droplets was proposed to exhibit a broad distribution of interatomic spacings if one considers the average overall interfacial regions of a nanoglassy specimen (Fig. 11). This speculation agrees with recent measurements by M6ssbauer spectroscopy

ULTRAFINEMICROSTRUCTURES

13

A

B

Fig. 11

Schematic cross sections through a two-dimensional nanoglass. The atoms are represented by circles. The material consists of small regions. In the central part of these regions (full circles), the interatomic spacings are similar to a bulk glass. In the interfacial areas broken lines, open circles), a broad spectrum of interatornic spacings may exist.

1.00 r-. .o

0.98

-~ 0.96 ,= 1.00 C:3 r'~

_~0.97 0.%

-2

-1 0 1 Velocity [mm/s]

2

0

1 OS [mm/s] l

Fig. 12a

M6ssbauer spectrum and distribution, p(QS), of the quadrupole splitting (QS) of a melt spun PdToSi27Fea metallic glass (47).

Fig. 12b

M6ssbau¢r spectrum and distribution, p(QS), of the quadrupole splitting (QS) of a nanoglass with the same chemical composition as in Fig. 12a. The diameter of the droplets used to generate the glass was about 4 nm. The quadrupole splitting consists of the following two components. A narrow one (solid line) similar to the quadrupole splitting of Fig. 12a and a broad component (broken line, hatched area) which is observed in the nanoglass only (47).

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(47). The observed M6ssbauer spectra of a PdToSizTFe3nanoglass and the corresponding quadrupole splitting distribution are compared in Fig. 12b with the M6ssbauer spectrum of a chemically identical glass prepared by melt spinning (Fig. 12a). Obviously both materials exhibit different spectra and different distributions of the quadrupole splitting. In fact, the quadrupole splitting of the nanogiass may be considered as a superposition of the following two components. One (narrow) component (solid line in Figs. 12a and 12b) which corresponds to a conventional (e.g., melt-spun) glass and a second (broad) component (broken line in Fig. 12b) which is observed in the nanoglass only. This result may be understood as follows. The quadrupole splitting depends on the nearest neighbor configuration around the iron atoms. Hence, the nanoglass consists structurally of two components, one component in which the atoms are arranged similarly to a melt-spun Pd-Si glass. In Fig. 11, this component is formed by the regions consisting of the atoms represented by the full circles. In the second component, the distribution of the interatomic configurations between nearest neighbors is much wider than in the melt-spun glass, as suggested by the broad distributions of the quadrupole splitting. This broad distribution is suggested to originate from the interdroplet regions (as was indicated in Fig. 11 by the open circles). Hence, the structure of nanogiasses may be manipulated in a controlled way as follows. The volume fraction of the interdroplet regions decreases if the droplets become larger. As a consequence, by varying the diameter and/or chemical composition of the consolidated glassy droplets, the relative volume fractions of the glassy and the broad second component can be manipulated. In other words, by varying the size of the giassy droplets, one may be able to vary (at constant chemical composition) the atomic structure of nanoglasses continuously between a melt-spun glass and a glassy structure with a broad distribution of interatomic spacings. As a matter of fact, this spectflation agrees with the experimental observation reported for Pd-Si-nanoglasses with different droplet diameters (47). If the chemical composition of the surface regions of an isolated glass droplet differ from the interior of the droplet (e.g., due to surface segregation effects), the interfacial regions deviate chemically as well as structurally from the center. Although the atomic structure of nanogiasses has been discussed in similar terms, as was done for nanocrystalline materials, there seems to be the following difference between both materials. Nanocrystalline materials preserve the low energy structure in the interior of the crystailites at the expense of the boundary regions, where the misfit between adjacent crystal lattices with different orientations is concentrated, so that a structure far away from equilibrium (with a large free volume and a broad distribution of interatomie spacing) is formed. In the case of nanogiasses, the glassy interior of the droplets (full circles in Fig. 11) is not a lowest energy structure. Hence, the atomic misfit of the boundary region between adjacent glassy droplets (e.g., between A and B in Fig. 11) in the as-consolidated state may not remain in this position if the nanogiass is held for some time at ambient temperature. In fact, the diffusionai rearrangement of individual atoms allows the system to redistribute the misfit and free volulne from the interfacial regions in the as-consolidated morphology to other regions within the structure in order to reduce the totai free energy of the systems. Hence, some time after consolidation, a nanoglass may have lost its "structural memory" to the asconsolidated morphology in the sense that the free volume distribution of the as-consolidated state is replaced by a new free volume distribution of lower free energy. In fact, such rearrangements may have been observed recently (50) by means of sinai1 angle scattering experiments on Au-Si nanoglasses. The proposed rearrangement also agrees with the result of computer simulations of glassy structures into which '~acancies" or "dislocations" have been introduced by removing an atom or by local shear. In other words, nanoglasses may exhibit structural (microstructure and atomic structure) features which differ from the ones of nanocrystalline materials, as well as from glasses prepared by conventional methods. Submicron Materials

In several recent papers (51) - (55) evidence has been presented suggesting that polycrystals (called submicron-grained materials) with grain sizes between 100 and 1000 nm exhibit properties deviating from those of coarse-grained polycrystais with the same chemical composition. For example, the Curie temperature of Ni decreased by 40K ff the grain size was reduced to 70 nm. Similarly, the diffusivity in Ni and Al was noticed to be enhanced by a factor of about 103 or more for grain sizes between 70 and 300 nm. These effects were interpreted by postulating that the regions between neighboring grains have different properties than the interior of the crystals. In fact, experimental evidence supporting this idea comes from studies by M6ssbauer

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spectroscopy of Fe with an average grain size of about 100 nm (55). The M6ssbaner spectrum consisted of two components, one of which agreed with the spectrum of bcc a-Fe, whereas the second component exhibited a reduced hyperfine field and an enhanced isomer shift. From the intensities of both components at low temperatures it was concluded that in polycrystalline Fe the boundary region, which differs propertywise from the crystal interior, is about 16 nm wide. In terms of the presently existing theories of the structure of grain boundaries (21), (22), it seems difficult to understand how a grain boundary can affect a region as wide as 8 nm. For example, measurements of the short circuit diffusion in isolated boundaries suggest a width between 0.5 to 1 nm for the region in which the diffusivity is enhanced (56). For practical applications, submicron grained materials seem to be attractive because they can be produced economically in large quantities by rapidly straining a material under pressure until a true strain of about 6 is achieved (53). Plasticity of Ceramics and lntermetallics Conventionally brittle ceramics have been observed to become ductile, permitting large (e.g., 100%) plastic deformations at low temperature (e.g., 293 K), if the ceramic is generated in the nanocrystalline form (57). The possibility to utilize this plasticity for net shape forming processes has been demonstrated recently at low temperatures where little or no grain growth occurs (58). If the plastic deformation is performed at temperatures of about 50% of the absolute melting point of the ceramic (e.g. for TiO2 deformation at about 800°C), total strains as high as 0.6 (without crack formation) were obtained in a period of about 15h (59). At 800°C, the grain size increased from about 50 nm to approximately 1/~m within 15h. During the deformation, the grains were noticed to remain equiaxed. With the exception of a few residual pores at grain triple junctions, no porosity was detected. Strain rates as high as 8-10-5s-1 were observed as true stresses of 52 MPa. These rates are quite sufficient for technological applications. The increasing grain size during deformation does not seem to pose a basic problem, in utilizing the ductility technologically. Doping of nanocrystalline ceramics (e.g. TiO2 by Y or AI) has been demonstrated to reduce grain growth dramatically. Grain boundary sliding, grain rotation and grain shape accommodation by diffusional processes seem to play a crucial role in the deformation of nanocrystalline ceramics, i.e., processes which are typical for superplastidty (60). In fact, superplastic flow in ceramics has been reported previously, however, at temperatures that were considerable higher and technologically more difficult to achieve. Creep with high strains was reported in AlzO3 at 1750-1950°C (61), AI203 doped with Cr2Oa and YzOa at 1500°C (62) and MgO doped AIzO3 at 1420°C (63). Only in ceramics with liquid phases at the interfaces was superplastic flow found at lower temperatures (64). Clearly, if the ductility of nanocrystalline ceramics is based on boundary sliding, grain rotation and accommodation, brittle inter-metallic compounds may also be expected to exhibit ductility in the nanocrystalline form. If this would be so, it may open the way to utilize intermetaUic compounds technologically. Their technological application has been hampered so far by brittleness. Ductile nanocrystalline ceramics may be used as a new type of ceramic material in their own fight. However, the plasticity may also be utilized for processing by extrusion or rolling. Subsequently, the material may be fully or partially converted back into a conventional ceramic. Partial conversion (e.g., by surface annealing) results in a material which exhibits at the surface the properties (e.g., hardness and chemical resistivity) of a conventional ceramic. However, the interior of such a material would still be ductile. The nanocrystalline approach deviates from the conventional strategy to defeat the brittleness of ceramics by improvements in processing and/or composition. Macrodefect-free cement, chemically modified silicon nitrides and transformation toughened zirconia (65), (66), represent successful improvements based on these two strategies. The last of these was originally proposed (67) with the provocative title "Ceramic steel?", because doped zirconia, like high-tensile steel, depends on a partial phase transformation to temper strength with shock resistance: an advancing crack triggers a local transformation which hinders further propagation of the crack. None of these improvements has gone far enough. The technological, economical and organizational problems that dimmed the promise of silicon nitride, which has now been the ceramic of the decade for three decades are described in ref. 68. The ceramic engine, outside Japan at least, still seems a long way from practical realization. Obviously, there is still a scope for new approaches.

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Magnetic Properties Measurements of the saturation magnetization (M=) of nanocrystalline iron (6nm crystal size) revealed a reduction of M, by about 40% relative to the saturation magnetization of bulk a-iron (35). For comparison, in metallic iron glasses (extrapolated to pure iron), M= is reduced by about 2% relative to a-Fe (35). The remarkable reduction of M° by 40% was attributed to the deviation of the interatomic spacings in the interracial regions as compared to the a-iron structure. This interpretation is supported by Mrssbauer spectroscopy measurements of the quadrupole splitting of nanocrystalline, anfiferromagnetic FeFz (69) indicating a wide distribution of different atomic configuration in the boundaries. If this is so, then one might speculate that other materials that are not ferromagnetic in the crystalline state may become ferromagnetic in the nanocrystalline form. Ferromagnetism in the boundary regions may not only result from local structural variations but also from an enhanced solute concentration. In fact, the solute solubility of various nanocrystalline systems such as H dissolved in Pd, Bi in Cu, Bi in Fe, Ag in Fe has been shown (43) (70), (71) to be enhanced by one or several orders of magnitude by incorporating the solute atoms in the cores of the boundaries or in the adjacent regions. Hence, in nanocrystalline alloys, the boundary regions differ structurally as well as chemically from the crystalline state, which may result in ferromagnetism in the boundary regions of nanocrystalline alloys, even if the crystalline state is not ferromagnetic. Recently, attractive magnetic properties have been reported for nanocrystalline Fe-alloys (72) generated by partially crystallizing an Fe-Si-B glass. The resulting material exhibited core losses as low as 200 mW/cm 3 and effective permeabilities of 105 or more. Some observations suggest that the ferromagnetic properties of the boundary regions play a critical role for these remarkable properties. Another feature which may be of relevance for future application of nanocrystaUine substances as storage materials, is their magnetic microstructure. The magnetic microstructure of nanocrystalline Fe was found to differ from the one of crystalline and amorphous Fe and iron alloys (73). The magnetic microstructure of crystalline and amorphous Fe consists of ferromagnetic domains separated by domain walls. In nanocrystalline Fe no domain structure was revealed by Bitter technique, Kerr microscopy and/or Lorentz electron microscopy (73). In fact, every crystaUite of a nanocrystalline Fe specimen seems to be a single ferromagnetic domain. The orientation of the magnetization of neighboring crystallites is controlled by two factors: the crystal anisotropy (which tries to align the magnetization of every crystallite in one of the easy directions) and the magnetic interaction between neighboring crystallites (which tries to align the magnetization of adjacent crystals into a common direction). If the crystallites are crystallographically oriented at random, a magnetic microstructure results, in which the magnetization of the crystallites is correlated over to a few crystal diameters (e.g. a few nanometers). Long range correlations resulting in domain formation are prevented by the random crystal orientation and the crystal anisotropy. The local magnetic structure can be changed by local magnetization on a nanometer scale suggesting the application as storage materials.

Perspectives One of the basic problems that has to be solved in order to permit the technological application of materials with nanometer-sized grains is the availability of economical methods to nroduce large auantities of such substances. The present situation seems to be somewhat analogous to metallic glasses. Shortly after the discovery of metallic glasses, several studies revealed technologically attractive properties of these materials. Nevertheless, it was the development of the melt spinning method (74), (75), which opened the way to the technological applications because it allowed to produce economically large quantities. Similarly, the present methods to generate nanocrystalline materials seem too expensive for large scale production. In fact, there are many chemical methods known for synthesizing economically large quantities of small crystals. Perhaps one of them can be utilized in the future. A second area which seems to deserve attention is the thermal stability of ultrafine microstructures. The high energy stored in these structures provides driving forces for recovery and reerystallization processes. However, ultrafine microstructures can be thermally stabilized in the form of multiphase materials consisting of crystals with different chemical compositions. If these crystals are mutually insoluble and if the various crystals are arranged in a nonperculative manner, thermally stable, ultrafine microstructures may be generated. So far, ultrafine microstructures formed by incorporating a high density of grain boundaries have been studied predominantly. Obviously, ultrafme microstructures by incorporating a high density of dislocations or stacking faults, etc. are likely to exhibit interesting properties as well (3).

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Acknowledgements This work has been supported by the BMFT under contract No. 523-400303 M 234 and the Deutsche Forsehungsgemeinschaft (G.W. Leibniz-Programm). The cooperation with colleagues from the Universities of Miinchen, Saarbrficken and Stuttgart, as well as from DESY Hamburg and ILL Grenoble, France, is appreciated. References 1. 2. 3. 4. 5. 6. 7. 8. 9. 10. 11. 12. 13. 14. 15. 16. 17. 18. 19. 20. 21. 22. 23. 24. 25. 26. 27. 28. 29. 30. 31.

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