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Amorphization by dislocation accumulation in shear bands
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Acta Materialia 57 (2009) 2851–2857 www.elsevier.com/locate/actamat
Amorphization by dislocation accumulation in shear bands Z.J. Lin a,b,1, M.J. Zhuo a,b, Z.Q. Sun a,b, P. Veyssie`re c, Y.C. Zhou a,* a
Shenyang National Laboratory for Materials Science, Institute of Metal Research, Chinese Academy of Sciences, 72 Wenhua Road, Shenyang 110016, China b Graduate School of Chinese Academy of Sciences, Beijing 100049, China c LEM, CNRS ONERA, 92322 Chaˆtillon Cedex, France Received 2 December 2008; received in revised form 16 February 2009; accepted 27 February 2009 Available online 1 April 2009
Abstract Microcrystalline c-Y2Si2O7 was indented at room temperature and the deformation microstructure was investigated by transmission electron microscopy in the vicinity of the indent. The volume directly beneath the indent comprises nanometer-sized grains delimited by an amorphous phase while dislocations dominate in the periphery either as dense slip bands in the border of the indent or, further away, as individual dislocations. The amorphous layers and the slip bands are a few nanometers thick. They lie along well-defined crystallographic planes. The microstructural organization is consistent with a stress-induced amorphization process whereby, under severe mechanical conditions, the crystal to amorphous transformation is mediated by slip bands containing a high density of dislocations. It is suggested that the damage tolerance of c-Y2Si2O7, which is exceptional for a ceramic material, benefits from this transformation. Ó 2009 Acta Materialia Inc. Published by Elsevier Ltd. All rights reserved. Keywords: Dislocation; Shear bands; TEM; Ceramics
1. Introduction Amorphization is a solid-state transformation that can be achieved in almost every category of crystalline materials including metal alloys [1–3], intermetallics [1,4,5], semiconductors [6–10], ceramics [11,12], minerals [13] and very few organic compounds [14]. The crystalline to amorphous (c ? a) transformation proceeds from the massive displacement of atoms into metastable positions and takes place when the crystal has reached a non-equilibrium state whose free energy is higher than that of the amorphous phase [5,15–17]. It is sometimes associated with a shear instability [18]. It can be achieved along several processing
*
Corresponding author. Address: Shenyang National Laboratory for Materials Science, Institute of Metal Research, Chinese Academy of Sciences, 72 Wenhua Road, Shenyang, 110016, China. Tel.: +86 24 2397 1765; fax: +86 24 2389 1320. E-mail address: [email protected] (Y.C. Zhou). 1 Present address: LANSCE-LC MS-H805, Los Alamos National Laboratory, NM 87545, USA.
routes such as irradiation [12,19], interdiffusion [3,20,21] or dehydration [14]. It can be also activated under an indenter [7,22,23] or upon shock-wave loading [11] and other severe mechanical conditions such as large deformation [24,25], mechanical alloying (MA) and mechanical milling (MM) of crystalline powders [1,2,5,6,8,17,26,27]. The mechanism involved in the c ? a transformation depends on the processing route. In MA, crystal destabilization is not purely a mechanical process but involves a solid-state reaction, whereas in MM, amorphization is thought to result in part from the accumulation of structural defects [17]. Whereas the nucleation of amorphization is not necessarily mediated by crystal defects as shown by Si [7,9,10] and B4C [11], there is repeated evidence of deformationinduced amorphous phases being nucleated at highly distorted lattice sites such as grain boundaries, twins, and dislocations [8,22]. Grain boundaries have a substantial and even overwhelming influence as this was shown by varying the size of nanoscale grains (for reviews, see Refs. [24,25]). The part taken by dislocations in deformation-induced amorphization is, on the other hand, rather ill-defined. In
1359-6454/$36.00 Ó 2009 Acta Materialia Inc. Published by Elsevier Ltd. All rights reserved. doi:10.1016/j.actamat.2009.02.040
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addition to being nucleation sites, dislocations are thought to favor the solid-state transformation in enhancing diffusion in the case of metal couples [28]. It has been shown that the dislocation density is dramatically increased under severe deformation conditions such as MA, where dislocations self-organize in shear bands forming cells, thus taking part in grain refinement [29], but there is no indication that the dynamical properties of dislocations such as self-organization as cells, storage or shear localization contribute directly to amorphization. The development of a dense dislocation structure during MM is sometimes reported [30] though not interpreted. In fact, transmission electron microscopy (TEM) observations have almost never been conducted at a scale adequate to pinpoint a direct relation between this and the nucleation of an amorphous phase. In the few documented instances of a c ? a transformation forming amorphous layers oriented in the host crystal such as in B4C [11], dislocations were not observed and it was thus concluded that they played no role. Quartz, on the other hand, transforms differently, depending upon whether it is compressed to high pressure in which case amorphization occurs along with shear, or else shocked, in which case it does not [31]. However, amorphization seems to precede shear band formation. In effect, it is by annealing amorphous lamellae, the dominant deformation structure in naturally shocked a-quartz, that planar deformation structures including dislocation bands, Brazil-twin lamellae and transformation lamellae are generated [13,32]. On the other hand, depending on load orientation, a-berlinite AlPO4, which is isostructural to a-quartz, deforms either by individual dislocations or by amorphous shear layers [31]. Under certain load orientations, pervasive amorphous shear lamellae are found in the weak planes of berlinite though no trace of dislocation activity is detected. In brief, while severe deformation normally involves dislocations in great numbers, no direct evidence on the role played collectively by dislocations in deformation-induced amorphization has been given and it is unclear whether there is a relation between mobile dislocation and the c ? a transformation. In this work, the defect distribution in the vicinity of indented c-Y2Si2O7 samples was investigated by transmission electron microscopy under diffraction contrast and high-resolution imaging. c-Y2Si2O7 is an interesting ceramic material in several respects. It is a grain boundary phase in Si3N4 ceramics doped with Y2O3 and/or SiO2 as sintering aids and, as such, it plays a beneficial role in enhancing the high-temperature mechanical properties of Si3N4 [33–35]. Y2Si2O7, which can be prepared as a single-phase bulk material [36], may be used in harsh environments under corrosive media or fast cooling/heating rates due to its good erosion resistance and low coefficient of thermal expansion [37]. With a melting point of 1775 °C, it is one of the most refractory silicates. Y2Si2O7 is therefore an excellent candidate as a high-temperature structural ceramic as well as an environmental/thermal barrier coating and oxidation protective coating on Si3N4 and SiC-
based composites. The mechanical properties of Y2Si2O7 are incompletely characterized, in particular the origin of its excellent damage tolerance and good machinability is not understood [38]. In the observations reported below, we investigate the deformation microstructure under indentation and we show that amorphous layers are formed along specific crystallographic planes as a result of considerable strains highly localized in dislocation slip bands. 2. Experimental c-Y2Si2O7 powders were synthesized through a solid/ liquid reaction method. The starting yttria and silica powders were calcined with 3 mol.% LiYO2 additive at 1400 °C for 4 h in air. Bulk polycrystalline c-Y2Si2O7 sample was prepared by pressureless sintering of the as-synthesized cY2Si2O7 powders at 1200 °C for 80 min in air. A detailed description of the synthesis process can be found in a previous publication [36]. c-Y2Si2O7 is the high-temperature stable phase among six different polymorphs (y, a, b, c, f, and g). c-Y2Si2O7 prepared by the present method can maintain the structural stability from room temperature to 1500 °C. c-Y2Si2O7 crystallizes as a monoclinic lattice with the P21/c space group [39]. It includes four Y2Si2O7 formulas per unit cell. Therefore the crystal structure can be described as a framework of YO6 octahedra and Si2O7 pyrosilicates alternately stacked in the three dimensions. The Si atoms are 4-coordinated to one O atom in a distorted tetrahedral environment while Y atoms are 6-coordinated to O, forming the YO6 octahedron. Two adjacent SiO4 tetrahedra share one O at the corner, forming the Si2O7 pyrosilicate structure with a linear Si–O–Si bridge (several projections of this lattice will be displayed in the following). Bulk Y2Si2O7 samples with dimensions of about 4 4 3 mm3 were indented by a spherical Si3N4 indenter with a diameter of 3.5 mm. The indenter was attached to a universal testing machine whose crosshead speed was 0.05 mm min1. No crack but a plastic zone was observed on the Y2Si2O7 sample surface, and thus the fracture toughness could not be deduced. The hardness of the as-prepared Y2Si2O7 ceramic is 6.2 GPa measured by Vickers indention, and fracture toughness is 2.12 MPa m1/2 determined by the chevron-notched beams method [38]. Samples for TEM analysis were prepared from regions containing indentations by polishing the face opposite to the indented surface, mounting the polished slices (10 lm thickness) on copper grids, and ion-milling at 4 kV. A 300 kV Tecnai G2 F30 TEM was used for both selected area electron diffraction (SAED) and high-resolution TEM (HRTEM) analyses. 3. Results Defect organization in indented c-Y2Si2O7 was investigated at various locations relative to the indent, hence in relation with the stress applied locally. Three microstructurally distinct zones could be distinguished according to
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the dominating microstructural feature. Fig. 1 shows isolated dislocations populating the outer region of the indented sample at a distance of approximately 20 lm away from the indent (in a volume subjected to relatively low stresses). The deformation microstructure, as viewed under the so-called weak-beam TEM imaging mode, is dominated by individual dislocations whose Burgers vectors and slip properties have not been characterized. The following is therefore, limited to general considerations. In Fig. 1a, the isolated dislocations that expand in their slip plane exhibit favored crystallographic orientations, suggesting a strong stabilizing interaction with the lattice, the so-called Peierls valleys. The dislocations are often dissociated into two partials with different Burgers vectors about 8 nm apart in projection (white arrow pairs in Fig. 1a and c). Dissociated dislocations sometimes exhibit constrictions with significant linear density (Fig. 1c), suggesting that they had been repeatedly intersected by others from other slip systems engendering constricted jogs (e.g., C) and hampering dislocation motion. An alternative explanation to the presence of constrictions (and jogs) in terms of repeated cross-slip is ruled out in the following paragraph. It is seen that in this area the isolated dislocations may coexist with planar arrays of glissile dislocations in places as in Fig. 1b and c (the letter H indicates a pile-up spearhead). The fact that dislocations are dissociated encourages planar slip and is consistent with the formation of such arrays and their piling up at obstacles. Roughly concentric with both the indent and the outer region, a second volume is distinguished from the frequency of its slip bands. This volume is submitted to stresses larger than those applied to the outer zone. Each slip band comprises dislocations with Burgers vectors all identical in direction and sign (see inset in Fig. 2a). They are grouped in bands under a density far larger than that commonly observed in the peripheral zone (Fig. 1a). The alignment of the dislocation extremities emerging at the free
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Fig. 2. Slip bands in the intermediate region of the indented volume. (a) Slip band planarity and evidence of profuse pile-ups. The dislocations exhibiting paired lines in the boxed area are not dipoles but dissociated dislocations since the distance between partials is constant whether the dipole is imaged with the g or g reflecting plane. (b) Intersecting slip bands. (c) The slight misalignment and differences in pile-up projected widths indicate that the slip bands are parallel to at least two crystallographically distinct planes.
surfaces of the thin foil indicates that dislocations are coplanar within ±3 nm, the resolution achieved under the present imaging conditions. The extreme planarity of slip in well-separated arrays is an indication that cross-slip is generally precluded for this would otherwise encourage slip
Fig. 1. Isolated dislocations populating the outer region of the indented sample, viewed under the so-called weak-beam TEM imaging mode. For enhanced visibility, the negative prints of the micrographs are displayed. (a) The isolated dislocations that expand in their easy slip plane. (b) A series of dislocations arranged to form pile-ups. (c) The dislocations show large linear densities of constrictions. The letter H indicates the pile-up spearhead. The images have the same scale bar.
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band thickening (Fig. 1a). In this region and at variance with what is encountered in the peripheral volume, isolated dislocations are scarce. Several slip systems are activated, delimitating cells of micrometric dimensions (Fig. 2b). A dislocation array may either be blocked by another one forming a Y-like or a T-like configuration, or it cross through as is manifested in places in the intersected bands (white arrows in Fig. 2b). Fig. 2 shows that dislocation arrays may delimitate elongated cells. The variability in cell shape takes its origin in the orientation of the grain explored relative to the applied stress field and the slip systems thus favored. The third, inner zone, which was in contact with the indenter tip and subjected to the highest stresses, displays about the same localized patterns as those found in the second zone with the distinction that the bands are amorphous. Fig. 3a displays an image realized with the signal scattered by the amorphous volume selected in between diffracted spots (e.g., white ring in bottom-left inset), the amorphous layers appears as bright lines within an otherwise almost featureless backgrounds. The bands exhibit two preferential orientations clearly visible here because they are seen reasonably edge on under this orientation of the foil. A third orientation can be distinguished, though with some difficulty, as an acute wedge-shaped contrast in the background, whose faintness is due to its inclination in the foil (letter V). Its wedge-shape reflects the foil decreasing thickness. The inset in bottom right of Fig. 3a is an image of the amorphous band taken in the high-resolution mode showing a band of amorphous material exhibiting an aperiodic signal about 6 nm wide surrounded by crystalline material on both sides with almost the same orientation (as revealed by the periodic projected potential image). The HREM image was taken with electron beam parallel to the [1 0 0] direction. The amorphous band is parallel to
the (0 0 1) plane. The image of another area obtained by selection of the transmitted beam is shown in Fig. 3b. The amorphous bands again appear as bright lines but large differences in crystal orientation are now manifested in the background by changes in intensity from dark (areas in Bragg orientation) to white (no signal diffracted). The thickness of the amorphous layers, which amounts to 6 ± 3 nm, is fairly constant over several hundreds nanometers (e.g., Fig. 3a, bottom-right inset). The coexistence of amorphous layers and pile-ups is hard to evidence in practice because the imaging conditions adequate to view amorphous layers are usually detrimental to pile-up visibility. The white arrows in Fig. 3b point at some of these pileups, manifested by changes in background contrast due to strong local distortions. The amorphous layers pertain to several main crystallographic habit planes: (0 1 0), (0 0 1), and within 10° from ð0 1 1Þ (Fig. 4). The latter non-crystallographic orientation results from the succession of amorphous slices in the two weakly bonded ð0 1 1Þ and ð0 2 1Þ planes which are only 18.6° apart and where the possibility of breaking their bonds and therefore amorphizing are almost the same. It should be noted that the slip bands such as those shown in Fig. 3 are in general elongated in directions parallel to those of the amorphous layers. It could not be ascertained though that all slip band orientations correspond to one of the above-mentioned habit planes of the amorphous layers. Layer intersections may assume either one of two configurations: (i) threefold junctions suggesting layer branching (letter Y in Fig. 5a) or (ii) mutually intersecting amorphous layers delimiting a pattern of sub-micron lozenge-like crystalline cells (Fig. 5). As for slip bands, one of the two intersecting amorphous bands is sometimes sheared whereas the other crosses through unaffected, indicating that the latter was sheared by the former (e.g., letter
Fig. 3. Distribution of amorphous bands in the inner volume of contact with the indenter, viewed under two complementary imaging modes. (a) Image taken with the signal scattered by the amorphous volume (as shown in bottom-left inset), the amorphous layers appears as bright lines within an otherwise featureless backgrounds. Displayed in the bottom-right inset is an HRTEM image of the amorphous band parallel to (0 0 1) showing the crystal potential projected along [1 0 0] and separated by a band of a periodic signal about 6 nm wide. (b) Bright-field image of another area. The amorphous bands again appear as bright lines but large differences in crystal orientation are manifested in the background by changes in intensity from dark (strongly diffracting) to white (weakly diffracting). The white arrows point regions of local distortion associated with dislocation pile-ups.
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of observation. The elongated cells are intersected by pileups in places. Amorphous layers are occasionally aligned with a pile-up, the two being separated by an oblique amorphous layer. Finally, adjacent grains exhibit slight misorientations that can be detected either by selected area diffraction or by direct inspection, through bend contours in neighboring subgrains (black arrowheads in Fig. 4). 4. Discussion
Fig. 4. A TEM image of amorphous layers forming in Y2Si2O7 with corresponding diffraction pattern shown in the inset. Elongated cells are presented. The diffraction pattern shows the three interfacial orientations favored under this projection. The layers are hardly seen but they can be deduced from the differences in contrast exhibited by neighboring crystalline cells. The boxed area contains slightly misoriented cells as indicated by the slight shifts between bend contours from one cell to the next (black arrowheads).
S in Fig. 5). In a number of cases though, intersections also occur with no apparent shift of either layer, which might result from the shear being mostly parallel to the direction
The similarities between the organization and properties of amorphous layers and those of slip bands in the inner and intermediate regions are striking. The morphology of the crystal cells is actually much the same whether they are delimited by slip bands in the intermediate volume or by amorphous layers in the inner region. In the latter, the two cell edges are, however, approximately 70 and 500 nm in length (Figs. 3a and 5a), that is, significantly less than the cell dimensions in the intermediate region. The volume subjected to the highest stresses during indentation has therefore undergone substantial grain refinement to near nanometer sizes, somewhat similar to observations in GaAs [22]. Some areas exhibit elongated, parallel amorphous layers, 150 nm in width on average and up to several microns in length (Figs. 3b and 4) with very few intersecting, oblique ones. This is again very similar to the elongated patterns encountered in the intermediate slipbanded region (Fig. 2a and c). Another important similarity between shear bands and amorphous layers lies in their crystallographic orientations which are generally identical as is reflected for example by the conspicuous alignments noticed in Fig. 3c. At variance from the highly covalent or iono-covalent Si, B4C, and SiO2, where dislocation movement is severely hindered, the bonding anisotropy of c-Y2Si2O7 results in low shear modulus and shear strength in the weakly
Fig. 5. Intersecting amorphous layers. (a) A general view showing regular lozenge-shaped crystalline cells coexisting with slightly elongated ones (bottom part). The inset is a magnified view of the boxed area illustrating the two branching modes, Y and S, between amorphous layers. At S, one of the two layers is sheared. (b and c) High-resolution images of the shear undergone by one layer upon intersection by another layer. This shear may reflect either the intersection of one slip band by another prior to amorphization, or the relative sliding of crystalline cells eased by the amorphous layers which are weaker than the crystalline matrix [41].
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bonded planes [27] and, in turn, in anisotropic mechanical properties. The Si–O bond, whose length varies between 0.1600 and 0.1616 nm, is much stronger than that between Y and O (the Y–O bond length lies between 0.2247 and 0.2288 nm). The projection of the crystal structure of c-Y2Si2O7 along different planes is shown in Fig. 6. The dashed lines highlight the unit cells of c-Y2Si2O7 while the parallel solid lines indicate the corresponding atomic planes. It is noted that the Y–O bonds dominate the bonding across these planes, consistent with the fact that the stress barrier to break the stronger Si–O bonds between tetrahedra is definitely higher than that to break the relatively weaker Y–O bonds. This is consistent with the present observations of preferred amorphous lying along the (0 0 1), (0 1 0), and ð0 1 1Þ and ð0 2 1Þ planes, with the latter two combined altogether to form the above-mentioned near ð0 1 1Þ layer. It is because of similar bonding configurations and crystallographic closeness (Fig. 6c and d) that the amorphous layers are seen lying near the ð0 1 1Þ or ð0 2 1Þ planes in the low magnification image (Fig. 4). One may envisage at least two ways for mobile dislocations to contribute to deformation-induced amorphization: (i) by forming a high supersaturation of point defects, i.e. interstitials and vacancies formed either by intersection or by mutual annihilation of dislocations, and/or (ii) by causing severe lattice distortions via exceptionally large dislocation densities. Both effects may of course operate concomitantly. Annihilation necessitates dislocations of both signs located at a few interplanar distances from each other in order for the dipole to annihilate spontaneously by forced climb at room temperature [40], which is clearly not the case here since slip bands are several tens to several hundreds of nanometers apart (Figs. 2 and 3). As for amorphization caused by dislocation accumulation, several factors are worth considering. One, inferred directly from the microstructure, is the apparent strength of the friction exerted by the lattice. It is well known that the stress ahead of a pile-up is that exerted by an individual dislocation times the number of dislocations in the pile-up. Hence, a
configuration such as that shown in Fig. 1b, which comprises about 20 visible dislocations piled up at H with no obstacle ahead, demonstrates that lattice friction alone is capable of withstanding shear stresses of the order of one to several GPa. From the inset in Fig. 2a, on the other hand, the mean spacing between consecutive dislocations is estimated to be 1.5 nm which, in a foil 200 nm thick, corresponds with a surface density qlocal = 1/(1.5 109 2 107) 3 1016 m2 or of the order of 1016 m2 when considered in a crystal layer several nanometers thick, that is, of the order of the thickness of an amorphous layer. This exceptionally high dislocation density has profound effects on strain and stress distribution in the pile-up vicinity and in particular on the hydrostatic stress. With respect to the hydrostatic stress at ±1 nm from an isolated dislocation, we find that the same stress component exerted by a planar array of 200 dislocations is more than doubled (the shear modulus and Poisson’s ratio were taken from [39] and assumed to be pressure independent). It is noticed that the frequent intersections between pile-ups and the subsequent dislocation accumulation (Fig. 2b) should also contribute to enhance local stresses. On the other hand, because of the jogs created on each dislocation by the other system (Fig. 1c), slip band intersection is also a source of point defects. Since the dislocations within a pile-up have all the same sign, the point defects are all of the same kind while it is difficult to envisage that the dislocations would simultaneously act as a source and as a sink for these point defects. The rate of defect production thus becomes locally higher than that of recovery engendering severely distorted volumes and raising the free energy of the system to a level higher than that of the amorphous phase so that it becomes possible for the amorphous phase to form by collapse of the topologically and chemically defected crystalline structure [41]. The internal stresses resulting from the accumulation of both point defects and dislocations can also relax by disordering more lattice sites near the amorphous regions thus promoting further expansion of the amorphous phase. Point defect reorganization within the pile-ups should be
Fig. 6. Schematic illustrations of weak-bonding planes within c-Y2Si2O7 (denoted by solid red lines and projected edge-on): (a) (0 1 0); (b) (1 0 0); (c) ð0 1 1Þ; and (d) ð0 2 1Þ. The dashed lines denote the unit cells. (a), (c), and (d) are projected along [1 0 0], whereas (b) is projected along [0 0 1].
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accelerated by pipe diffusion, an argument already invoked for amorphization by interdiffusion [28]. 5. Conclusion Extensive TEM analyses on the indented polycrystalline c-Y2Si2O7 ceramic were conducted. The similarities between the organization and properties of amorphous layers and those of slip bands in the inner and intermediate regions, respectively, reflect the specific mechanism involved in the c ? a transformation of c-Y2Si2O7, that is, the accumulation of planar arrays of dislocations building up localized lattice distortions. How and to what extent lattice defects might contribute to the growth of the amorphous phase in indented c-Y2Si2O7, and what implications these effects have on damage tolerance, can be outlined as follows. In c-Y2Si2O7, two factors appear to encourage the c ? a transformation. One is the absence of dislocation cross-slip resulting in planar slip, the other is a considerable lattice friction which enhances the planar dislocation density up to the phase transformation level. The higher the stress, the more pervasive the pile-ups, which explains the observed grain refinement in the near vicinity of the indenter. The presence and the organization of amorphous layers contribute to the exceptional damage tolerance of cY2Si2O7 for two main reasons. As discussed in Ref. [41], the amorphous layers ease the relative sliding of submicrometer crystals cells. Cracks nucleate wherever the sliding causes too large a strain incompatibility but because of the profusion of closely spaced amorphous layers, crack propagation is blunted within a fraction of a micrometer. Acknowledgements P. Veyssie`re wishes to thank Drs. P. Cordier, A. Finel, and Y. Le Bouar for discussions and the Hsun Lee foundation for supporting his visit at the Institute of Metal Research, Chinese Academy of Sciences. Z.J. Lin appreciates Profs. J.Y. Wang and M.S. Li for kind help. This work was supported by the National Outstanding Young Scientist Foundation for Y.C. Zhou (59925208) and Natural Sciences Foundation of China (50772114). References [1] Hellstern E, Schultz L. Appl Phys Lett 1986;18:124.
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Acta Materialia 57 (2009) 2851–2857 www.elsevier.com/locate/actamat
Amorphization by dislocation accumulation in shear bands Z.J. Lin a,b,1, M.J. Zhuo a,b, Z.Q. Sun a,b, P. Veyssie`re c, Y.C. Zhou a,* a
Shenyang National Laboratory for Materials Science, Institute of Metal Research, Chinese Academy of Sciences, 72 Wenhua Road, Shenyang 110016, China b Graduate School of Chinese Academy of Sciences, Beijing 100049, China c LEM, CNRS ONERA, 92322 Chaˆtillon Cedex, France Received 2 December 2008; received in revised form 16 February 2009; accepted 27 February 2009 Available online 1 April 2009
Abstract Microcrystalline c-Y2Si2O7 was indented at room temperature and the deformation microstructure was investigated by transmission electron microscopy in the vicinity of the indent. The volume directly beneath the indent comprises nanometer-sized grains delimited by an amorphous phase while dislocations dominate in the periphery either as dense slip bands in the border of the indent or, further away, as individual dislocations. The amorphous layers and the slip bands are a few nanometers thick. They lie along well-defined crystallographic planes. The microstructural organization is consistent with a stress-induced amorphization process whereby, under severe mechanical conditions, the crystal to amorphous transformation is mediated by slip bands containing a high density of dislocations. It is suggested that the damage tolerance of c-Y2Si2O7, which is exceptional for a ceramic material, benefits from this transformation. Ó 2009 Acta Materialia Inc. Published by Elsevier Ltd. All rights reserved. Keywords: Dislocation; Shear bands; TEM; Ceramics
1. Introduction Amorphization is a solid-state transformation that can be achieved in almost every category of crystalline materials including metal alloys [1–3], intermetallics [1,4,5], semiconductors [6–10], ceramics [11,12], minerals [13] and very few organic compounds [14]. The crystalline to amorphous (c ? a) transformation proceeds from the massive displacement of atoms into metastable positions and takes place when the crystal has reached a non-equilibrium state whose free energy is higher than that of the amorphous phase [5,15–17]. It is sometimes associated with a shear instability [18]. It can be achieved along several processing
*
Corresponding author. Address: Shenyang National Laboratory for Materials Science, Institute of Metal Research, Chinese Academy of Sciences, 72 Wenhua Road, Shenyang, 110016, China. Tel.: +86 24 2397 1765; fax: +86 24 2389 1320. E-mail address: [email protected] (Y.C. Zhou). 1 Present address: LANSCE-LC MS-H805, Los Alamos National Laboratory, NM 87545, USA.
routes such as irradiation [12,19], interdiffusion [3,20,21] or dehydration [14]. It can be also activated under an indenter [7,22,23] or upon shock-wave loading [11] and other severe mechanical conditions such as large deformation [24,25], mechanical alloying (MA) and mechanical milling (MM) of crystalline powders [1,2,5,6,8,17,26,27]. The mechanism involved in the c ? a transformation depends on the processing route. In MA, crystal destabilization is not purely a mechanical process but involves a solid-state reaction, whereas in MM, amorphization is thought to result in part from the accumulation of structural defects [17]. Whereas the nucleation of amorphization is not necessarily mediated by crystal defects as shown by Si [7,9,10] and B4C [11], there is repeated evidence of deformationinduced amorphous phases being nucleated at highly distorted lattice sites such as grain boundaries, twins, and dislocations [8,22]. Grain boundaries have a substantial and even overwhelming influence as this was shown by varying the size of nanoscale grains (for reviews, see Refs. [24,25]). The part taken by dislocations in deformation-induced amorphization is, on the other hand, rather ill-defined. In
1359-6454/$36.00 Ó 2009 Acta Materialia Inc. Published by Elsevier Ltd. All rights reserved. doi:10.1016/j.actamat.2009.02.040
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addition to being nucleation sites, dislocations are thought to favor the solid-state transformation in enhancing diffusion in the case of metal couples [28]. It has been shown that the dislocation density is dramatically increased under severe deformation conditions such as MA, where dislocations self-organize in shear bands forming cells, thus taking part in grain refinement [29], but there is no indication that the dynamical properties of dislocations such as self-organization as cells, storage or shear localization contribute directly to amorphization. The development of a dense dislocation structure during MM is sometimes reported [30] though not interpreted. In fact, transmission electron microscopy (TEM) observations have almost never been conducted at a scale adequate to pinpoint a direct relation between this and the nucleation of an amorphous phase. In the few documented instances of a c ? a transformation forming amorphous layers oriented in the host crystal such as in B4C [11], dislocations were not observed and it was thus concluded that they played no role. Quartz, on the other hand, transforms differently, depending upon whether it is compressed to high pressure in which case amorphization occurs along with shear, or else shocked, in which case it does not [31]. However, amorphization seems to precede shear band formation. In effect, it is by annealing amorphous lamellae, the dominant deformation structure in naturally shocked a-quartz, that planar deformation structures including dislocation bands, Brazil-twin lamellae and transformation lamellae are generated [13,32]. On the other hand, depending on load orientation, a-berlinite AlPO4, which is isostructural to a-quartz, deforms either by individual dislocations or by amorphous shear layers [31]. Under certain load orientations, pervasive amorphous shear lamellae are found in the weak planes of berlinite though no trace of dislocation activity is detected. In brief, while severe deformation normally involves dislocations in great numbers, no direct evidence on the role played collectively by dislocations in deformation-induced amorphization has been given and it is unclear whether there is a relation between mobile dislocation and the c ? a transformation. In this work, the defect distribution in the vicinity of indented c-Y2Si2O7 samples was investigated by transmission electron microscopy under diffraction contrast and high-resolution imaging. c-Y2Si2O7 is an interesting ceramic material in several respects. It is a grain boundary phase in Si3N4 ceramics doped with Y2O3 and/or SiO2 as sintering aids and, as such, it plays a beneficial role in enhancing the high-temperature mechanical properties of Si3N4 [33–35]. Y2Si2O7, which can be prepared as a single-phase bulk material [36], may be used in harsh environments under corrosive media or fast cooling/heating rates due to its good erosion resistance and low coefficient of thermal expansion [37]. With a melting point of 1775 °C, it is one of the most refractory silicates. Y2Si2O7 is therefore an excellent candidate as a high-temperature structural ceramic as well as an environmental/thermal barrier coating and oxidation protective coating on Si3N4 and SiC-
based composites. The mechanical properties of Y2Si2O7 are incompletely characterized, in particular the origin of its excellent damage tolerance and good machinability is not understood [38]. In the observations reported below, we investigate the deformation microstructure under indentation and we show that amorphous layers are formed along specific crystallographic planes as a result of considerable strains highly localized in dislocation slip bands. 2. Experimental c-Y2Si2O7 powders were synthesized through a solid/ liquid reaction method. The starting yttria and silica powders were calcined with 3 mol.% LiYO2 additive at 1400 °C for 4 h in air. Bulk polycrystalline c-Y2Si2O7 sample was prepared by pressureless sintering of the as-synthesized cY2Si2O7 powders at 1200 °C for 80 min in air. A detailed description of the synthesis process can be found in a previous publication [36]. c-Y2Si2O7 is the high-temperature stable phase among six different polymorphs (y, a, b, c, f, and g). c-Y2Si2O7 prepared by the present method can maintain the structural stability from room temperature to 1500 °C. c-Y2Si2O7 crystallizes as a monoclinic lattice with the P21/c space group [39]. It includes four Y2Si2O7 formulas per unit cell. Therefore the crystal structure can be described as a framework of YO6 octahedra and Si2O7 pyrosilicates alternately stacked in the three dimensions. The Si atoms are 4-coordinated to one O atom in a distorted tetrahedral environment while Y atoms are 6-coordinated to O, forming the YO6 octahedron. Two adjacent SiO4 tetrahedra share one O at the corner, forming the Si2O7 pyrosilicate structure with a linear Si–O–Si bridge (several projections of this lattice will be displayed in the following). Bulk Y2Si2O7 samples with dimensions of about 4 4 3 mm3 were indented by a spherical Si3N4 indenter with a diameter of 3.5 mm. The indenter was attached to a universal testing machine whose crosshead speed was 0.05 mm min1. No crack but a plastic zone was observed on the Y2Si2O7 sample surface, and thus the fracture toughness could not be deduced. The hardness of the as-prepared Y2Si2O7 ceramic is 6.2 GPa measured by Vickers indention, and fracture toughness is 2.12 MPa m1/2 determined by the chevron-notched beams method [38]. Samples for TEM analysis were prepared from regions containing indentations by polishing the face opposite to the indented surface, mounting the polished slices (10 lm thickness) on copper grids, and ion-milling at 4 kV. A 300 kV Tecnai G2 F30 TEM was used for both selected area electron diffraction (SAED) and high-resolution TEM (HRTEM) analyses. 3. Results Defect organization in indented c-Y2Si2O7 was investigated at various locations relative to the indent, hence in relation with the stress applied locally. Three microstructurally distinct zones could be distinguished according to
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the dominating microstructural feature. Fig. 1 shows isolated dislocations populating the outer region of the indented sample at a distance of approximately 20 lm away from the indent (in a volume subjected to relatively low stresses). The deformation microstructure, as viewed under the so-called weak-beam TEM imaging mode, is dominated by individual dislocations whose Burgers vectors and slip properties have not been characterized. The following is therefore, limited to general considerations. In Fig. 1a, the isolated dislocations that expand in their slip plane exhibit favored crystallographic orientations, suggesting a strong stabilizing interaction with the lattice, the so-called Peierls valleys. The dislocations are often dissociated into two partials with different Burgers vectors about 8 nm apart in projection (white arrow pairs in Fig. 1a and c). Dissociated dislocations sometimes exhibit constrictions with significant linear density (Fig. 1c), suggesting that they had been repeatedly intersected by others from other slip systems engendering constricted jogs (e.g., C) and hampering dislocation motion. An alternative explanation to the presence of constrictions (and jogs) in terms of repeated cross-slip is ruled out in the following paragraph. It is seen that in this area the isolated dislocations may coexist with planar arrays of glissile dislocations in places as in Fig. 1b and c (the letter H indicates a pile-up spearhead). The fact that dislocations are dissociated encourages planar slip and is consistent with the formation of such arrays and their piling up at obstacles. Roughly concentric with both the indent and the outer region, a second volume is distinguished from the frequency of its slip bands. This volume is submitted to stresses larger than those applied to the outer zone. Each slip band comprises dislocations with Burgers vectors all identical in direction and sign (see inset in Fig. 2a). They are grouped in bands under a density far larger than that commonly observed in the peripheral zone (Fig. 1a). The alignment of the dislocation extremities emerging at the free
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Fig. 2. Slip bands in the intermediate region of the indented volume. (a) Slip band planarity and evidence of profuse pile-ups. The dislocations exhibiting paired lines in the boxed area are not dipoles but dissociated dislocations since the distance between partials is constant whether the dipole is imaged with the g or g reflecting plane. (b) Intersecting slip bands. (c) The slight misalignment and differences in pile-up projected widths indicate that the slip bands are parallel to at least two crystallographically distinct planes.
surfaces of the thin foil indicates that dislocations are coplanar within ±3 nm, the resolution achieved under the present imaging conditions. The extreme planarity of slip in well-separated arrays is an indication that cross-slip is generally precluded for this would otherwise encourage slip
Fig. 1. Isolated dislocations populating the outer region of the indented sample, viewed under the so-called weak-beam TEM imaging mode. For enhanced visibility, the negative prints of the micrographs are displayed. (a) The isolated dislocations that expand in their easy slip plane. (b) A series of dislocations arranged to form pile-ups. (c) The dislocations show large linear densities of constrictions. The letter H indicates the pile-up spearhead. The images have the same scale bar.
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band thickening (Fig. 1a). In this region and at variance with what is encountered in the peripheral volume, isolated dislocations are scarce. Several slip systems are activated, delimitating cells of micrometric dimensions (Fig. 2b). A dislocation array may either be blocked by another one forming a Y-like or a T-like configuration, or it cross through as is manifested in places in the intersected bands (white arrows in Fig. 2b). Fig. 2 shows that dislocation arrays may delimitate elongated cells. The variability in cell shape takes its origin in the orientation of the grain explored relative to the applied stress field and the slip systems thus favored. The third, inner zone, which was in contact with the indenter tip and subjected to the highest stresses, displays about the same localized patterns as those found in the second zone with the distinction that the bands are amorphous. Fig. 3a displays an image realized with the signal scattered by the amorphous volume selected in between diffracted spots (e.g., white ring in bottom-left inset), the amorphous layers appears as bright lines within an otherwise almost featureless backgrounds. The bands exhibit two preferential orientations clearly visible here because they are seen reasonably edge on under this orientation of the foil. A third orientation can be distinguished, though with some difficulty, as an acute wedge-shaped contrast in the background, whose faintness is due to its inclination in the foil (letter V). Its wedge-shape reflects the foil decreasing thickness. The inset in bottom right of Fig. 3a is an image of the amorphous band taken in the high-resolution mode showing a band of amorphous material exhibiting an aperiodic signal about 6 nm wide surrounded by crystalline material on both sides with almost the same orientation (as revealed by the periodic projected potential image). The HREM image was taken with electron beam parallel to the [1 0 0] direction. The amorphous band is parallel to
the (0 0 1) plane. The image of another area obtained by selection of the transmitted beam is shown in Fig. 3b. The amorphous bands again appear as bright lines but large differences in crystal orientation are now manifested in the background by changes in intensity from dark (areas in Bragg orientation) to white (no signal diffracted). The thickness of the amorphous layers, which amounts to 6 ± 3 nm, is fairly constant over several hundreds nanometers (e.g., Fig. 3a, bottom-right inset). The coexistence of amorphous layers and pile-ups is hard to evidence in practice because the imaging conditions adequate to view amorphous layers are usually detrimental to pile-up visibility. The white arrows in Fig. 3b point at some of these pileups, manifested by changes in background contrast due to strong local distortions. The amorphous layers pertain to several main crystallographic habit planes: (0 1 0), (0 0 1), and within 10° from ð0 1 1Þ (Fig. 4). The latter non-crystallographic orientation results from the succession of amorphous slices in the two weakly bonded ð0 1 1Þ and ð0 2 1Þ planes which are only 18.6° apart and where the possibility of breaking their bonds and therefore amorphizing are almost the same. It should be noted that the slip bands such as those shown in Fig. 3 are in general elongated in directions parallel to those of the amorphous layers. It could not be ascertained though that all slip band orientations correspond to one of the above-mentioned habit planes of the amorphous layers. Layer intersections may assume either one of two configurations: (i) threefold junctions suggesting layer branching (letter Y in Fig. 5a) or (ii) mutually intersecting amorphous layers delimiting a pattern of sub-micron lozenge-like crystalline cells (Fig. 5). As for slip bands, one of the two intersecting amorphous bands is sometimes sheared whereas the other crosses through unaffected, indicating that the latter was sheared by the former (e.g., letter
Fig. 3. Distribution of amorphous bands in the inner volume of contact with the indenter, viewed under two complementary imaging modes. (a) Image taken with the signal scattered by the amorphous volume (as shown in bottom-left inset), the amorphous layers appears as bright lines within an otherwise featureless backgrounds. Displayed in the bottom-right inset is an HRTEM image of the amorphous band parallel to (0 0 1) showing the crystal potential projected along [1 0 0] and separated by a band of a periodic signal about 6 nm wide. (b) Bright-field image of another area. The amorphous bands again appear as bright lines but large differences in crystal orientation are manifested in the background by changes in intensity from dark (strongly diffracting) to white (weakly diffracting). The white arrows point regions of local distortion associated with dislocation pile-ups.
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of observation. The elongated cells are intersected by pileups in places. Amorphous layers are occasionally aligned with a pile-up, the two being separated by an oblique amorphous layer. Finally, adjacent grains exhibit slight misorientations that can be detected either by selected area diffraction or by direct inspection, through bend contours in neighboring subgrains (black arrowheads in Fig. 4). 4. Discussion
Fig. 4. A TEM image of amorphous layers forming in Y2Si2O7 with corresponding diffraction pattern shown in the inset. Elongated cells are presented. The diffraction pattern shows the three interfacial orientations favored under this projection. The layers are hardly seen but they can be deduced from the differences in contrast exhibited by neighboring crystalline cells. The boxed area contains slightly misoriented cells as indicated by the slight shifts between bend contours from one cell to the next (black arrowheads).
S in Fig. 5). In a number of cases though, intersections also occur with no apparent shift of either layer, which might result from the shear being mostly parallel to the direction
The similarities between the organization and properties of amorphous layers and those of slip bands in the inner and intermediate regions are striking. The morphology of the crystal cells is actually much the same whether they are delimited by slip bands in the intermediate volume or by amorphous layers in the inner region. In the latter, the two cell edges are, however, approximately 70 and 500 nm in length (Figs. 3a and 5a), that is, significantly less than the cell dimensions in the intermediate region. The volume subjected to the highest stresses during indentation has therefore undergone substantial grain refinement to near nanometer sizes, somewhat similar to observations in GaAs [22]. Some areas exhibit elongated, parallel amorphous layers, 150 nm in width on average and up to several microns in length (Figs. 3b and 4) with very few intersecting, oblique ones. This is again very similar to the elongated patterns encountered in the intermediate slipbanded region (Fig. 2a and c). Another important similarity between shear bands and amorphous layers lies in their crystallographic orientations which are generally identical as is reflected for example by the conspicuous alignments noticed in Fig. 3c. At variance from the highly covalent or iono-covalent Si, B4C, and SiO2, where dislocation movement is severely hindered, the bonding anisotropy of c-Y2Si2O7 results in low shear modulus and shear strength in the weakly
Fig. 5. Intersecting amorphous layers. (a) A general view showing regular lozenge-shaped crystalline cells coexisting with slightly elongated ones (bottom part). The inset is a magnified view of the boxed area illustrating the two branching modes, Y and S, between amorphous layers. At S, one of the two layers is sheared. (b and c) High-resolution images of the shear undergone by one layer upon intersection by another layer. This shear may reflect either the intersection of one slip band by another prior to amorphization, or the relative sliding of crystalline cells eased by the amorphous layers which are weaker than the crystalline matrix [41].
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bonded planes [27] and, in turn, in anisotropic mechanical properties. The Si–O bond, whose length varies between 0.1600 and 0.1616 nm, is much stronger than that between Y and O (the Y–O bond length lies between 0.2247 and 0.2288 nm). The projection of the crystal structure of c-Y2Si2O7 along different planes is shown in Fig. 6. The dashed lines highlight the unit cells of c-Y2Si2O7 while the parallel solid lines indicate the corresponding atomic planes. It is noted that the Y–O bonds dominate the bonding across these planes, consistent with the fact that the stress barrier to break the stronger Si–O bonds between tetrahedra is definitely higher than that to break the relatively weaker Y–O bonds. This is consistent with the present observations of preferred amorphous lying along the (0 0 1), (0 1 0), and ð0 1 1Þ and ð0 2 1Þ planes, with the latter two combined altogether to form the above-mentioned near ð0 1 1Þ layer. It is because of similar bonding configurations and crystallographic closeness (Fig. 6c and d) that the amorphous layers are seen lying near the ð0 1 1Þ or ð0 2 1Þ planes in the low magnification image (Fig. 4). One may envisage at least two ways for mobile dislocations to contribute to deformation-induced amorphization: (i) by forming a high supersaturation of point defects, i.e. interstitials and vacancies formed either by intersection or by mutual annihilation of dislocations, and/or (ii) by causing severe lattice distortions via exceptionally large dislocation densities. Both effects may of course operate concomitantly. Annihilation necessitates dislocations of both signs located at a few interplanar distances from each other in order for the dipole to annihilate spontaneously by forced climb at room temperature [40], which is clearly not the case here since slip bands are several tens to several hundreds of nanometers apart (Figs. 2 and 3). As for amorphization caused by dislocation accumulation, several factors are worth considering. One, inferred directly from the microstructure, is the apparent strength of the friction exerted by the lattice. It is well known that the stress ahead of a pile-up is that exerted by an individual dislocation times the number of dislocations in the pile-up. Hence, a
configuration such as that shown in Fig. 1b, which comprises about 20 visible dislocations piled up at H with no obstacle ahead, demonstrates that lattice friction alone is capable of withstanding shear stresses of the order of one to several GPa. From the inset in Fig. 2a, on the other hand, the mean spacing between consecutive dislocations is estimated to be 1.5 nm which, in a foil 200 nm thick, corresponds with a surface density qlocal = 1/(1.5 109 2 107) 3 1016 m2 or of the order of 1016 m2 when considered in a crystal layer several nanometers thick, that is, of the order of the thickness of an amorphous layer. This exceptionally high dislocation density has profound effects on strain and stress distribution in the pile-up vicinity and in particular on the hydrostatic stress. With respect to the hydrostatic stress at ±1 nm from an isolated dislocation, we find that the same stress component exerted by a planar array of 200 dislocations is more than doubled (the shear modulus and Poisson’s ratio were taken from [39] and assumed to be pressure independent). It is noticed that the frequent intersections between pile-ups and the subsequent dislocation accumulation (Fig. 2b) should also contribute to enhance local stresses. On the other hand, because of the jogs created on each dislocation by the other system (Fig. 1c), slip band intersection is also a source of point defects. Since the dislocations within a pile-up have all the same sign, the point defects are all of the same kind while it is difficult to envisage that the dislocations would simultaneously act as a source and as a sink for these point defects. The rate of defect production thus becomes locally higher than that of recovery engendering severely distorted volumes and raising the free energy of the system to a level higher than that of the amorphous phase so that it becomes possible for the amorphous phase to form by collapse of the topologically and chemically defected crystalline structure [41]. The internal stresses resulting from the accumulation of both point defects and dislocations can also relax by disordering more lattice sites near the amorphous regions thus promoting further expansion of the amorphous phase. Point defect reorganization within the pile-ups should be
Fig. 6. Schematic illustrations of weak-bonding planes within c-Y2Si2O7 (denoted by solid red lines and projected edge-on): (a) (0 1 0); (b) (1 0 0); (c) ð0 1 1Þ; and (d) ð0 2 1Þ. The dashed lines denote the unit cells. (a), (c), and (d) are projected along [1 0 0], whereas (b) is projected along [0 0 1].
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accelerated by pipe diffusion, an argument already invoked for amorphization by interdiffusion [28]. 5. Conclusion Extensive TEM analyses on the indented polycrystalline c-Y2Si2O7 ceramic were conducted. The similarities between the organization and properties of amorphous layers and those of slip bands in the inner and intermediate regions, respectively, reflect the specific mechanism involved in the c ? a transformation of c-Y2Si2O7, that is, the accumulation of planar arrays of dislocations building up localized lattice distortions. How and to what extent lattice defects might contribute to the growth of the amorphous phase in indented c-Y2Si2O7, and what implications these effects have on damage tolerance, can be outlined as follows. In c-Y2Si2O7, two factors appear to encourage the c ? a transformation. One is the absence of dislocation cross-slip resulting in planar slip, the other is a considerable lattice friction which enhances the planar dislocation density up to the phase transformation level. The higher the stress, the more pervasive the pile-ups, which explains the observed grain refinement in the near vicinity of the indenter. The presence and the organization of amorphous layers contribute to the exceptional damage tolerance of cY2Si2O7 for two main reasons. As discussed in Ref. [41], the amorphous layers ease the relative sliding of submicrometer crystals cells. Cracks nucleate wherever the sliding causes too large a strain incompatibility but because of the profusion of closely spaced amorphous layers, crack propagation is blunted within a fraction of a micrometer. Acknowledgements P. Veyssie`re wishes to thank Drs. P. Cordier, A. Finel, and Y. Le Bouar for discussions and the Hsun Lee foundation for supporting his visit at the Institute of Metal Research, Chinese Academy of Sciences. Z.J. Lin appreciates Profs. J.Y. Wang and M.S. Li for kind help. This work was supported by the National Outstanding Young Scientist Foundation for Y.C. Zhou (59925208) and Natural Sciences Foundation of China (50772114). References [1] Hellstern E, Schultz L. Appl Phys Lett 1986;18:124.
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