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Crystallographic assembly of macroscopic crystals by sub-parallel splicing of multiple seeds Victor G. Thomas, Nina Daneu, Aleksander Recnik, Dmitry A. Fursenko, Sergey P. Demin, Svjatoslav P. Belinsky, and Pavel N. Gavryushkin Cryst. Growth Des., Just Accepted Manuscript • DOI: 10.1021/acs.cgd.6b01616 • Publication Date (Web): 04 Jan 2017 Downloaded from http://pubs.acs.org on January 10, 2017

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Crystal Growth & Design

Crystallographic assembly of macroscopic crystals by sub-parallel splicing of multiple seeds Victor G. Thomas a*, Nina Daneu b, Aleksander Rečnik b, Dmitry A. Fursenko a, Sergey P. Demin с, Svjatoslav P. Belinsky d, Pavel N. Gavryushkin a,d a

Sobolev Institute of Geology and Mineralogy Siberian Branch Russian Academy of Science, pr. Academician Koptyug, 3, Novosibirsk 630090, Russia. b Department for Nanostructured Materials, Jožef Stefan Institute, Jamova cesta 39, 1000 Ljubljana, Slovenia. c Tairus Company, pr. Academician Koptyug, 3/4, Novosibirsk 630090, Russia. d Novosibirsk State University, Pirogova st., 2, Novosibirsk 630090, Russia. *

Corresponding author, e-mail: [email protected], [email protected].

Abstract The potential possibility of intergrowth of two bulk crystals to single crystal was demonstrated on the example of beryl, Be3Al2Si6O18, growing under hydrothermal conditions. The result has practical importance because it allows to increase the yield of useful product on a single growth cycle. The fine structure and composition of the matter of the area adjacent to the splicing boundary was investigated. It was demonstrated that the influence of the intergrowths border of two crystals spliced in parallel is entirely analogous to the affect of a twin boundary. This analogy extends as well on the specific morphology of the growth front generated by the boundary as on the growth velocity of the surfaces adjacent to the boundary that is increasing in 3÷10 times. In addition we show that the spliced crystals orientate each other in parallel. The assumption on the nature of the driving forces of such impact was made. We also suggest assumption about the nature of this orientational forces.

Introduction Hydrothermal (HT) growth is currently the only method used for industrial production of gem quality emerald crystals: i.e. beryl (Be3Al2Si6O18, space group P6/mcc), doped with Cr3+ or V3+ ions. Common to all HT approaches regardless the diversity in chemical composition of starting solutions or the source of growth material (either reactive synthesis from the primary oxides or recrystallization of natural beryl charge)1,2 is that they all necessitate a high-quality seed crystal for growth of flawless gem material. To achieve maximum growth rate for the appropriate quality of the material, the crystals are cut along some planes nonparallel to the main crystal faces 4, 8–13. Spe_ cifically, all emerald crystals, with rare exceptions3, are grown on seed plates cut in the zone [ 1 1 00] _ at ~30±3° angle with respect to the prismatic face a__{1120} . The approximate orientation of these seed plates can be expressed in Miller indices {5.5.10.6} ♥. This orientation of the seed plates provides a maximum growth rate for crystals with satisfactory quality 4–6. __

The surfaces, such as {5.5.10.6} , which are non-parallel to any of the possible faces of the beryl crystals7, are growing by the sawtooth-like front. We call them as regeneration surfaces8. ♥

Indexing is given relative to the unit cell metrics. Another common alternative is determining the indices with respect _ to the single face s {1121} – the secondary hexagonal bipyramid 7. In that case the Miller index l is two times smaller.

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__

Figure 1. Emerald crystal grown on seed plate with R {5.5.10 .6} regeneration surface, using _ _ HT method. _ The letters indicate faces, representing crystallographic forms c {0001} , m {1010} , a {11__20} , p {1012} _ and s {1122} . The macroscopically rough face on the top is growth front of the R {5.5.10 .6} surface.

Morphological aspects of how the regeneration surfaces grow on beryl crystals have been discussed in 4, 9, 10; while general theoretical overview of regeneration __growth process is given in 8,11,12. An example of emerald crystal grown on a seed parallel to R {5.5.10.6} is illustrated in Figure 1. The general shape of such __a crystal is_ _mainly defined by the rough macroscopic surfaces – the _ _ regeneration surfaces R (5.5.10.6) and (5.5.10.6) ; primary prism m faces with _indices (_1_100) and _ along seed length; and secondary prism faces a (1120) and (1120) and (1100) that are oriented _ _ _ the _ bipyramidal s (1122) and (1122) faces, tilted at angles of ~30° (both a) and ~15° (both s) with respect to the seed R-plane. It has been previously shown 4,13 that during recrystallization of natural beryl crystals 2,5, the average growth velocities, V, of corresponding faces are in the following ratios: VR : Va : Vs : Vp : Vm : Vc ≈ 100 : 15 : 12 : 6 : 6 : 1. The useful parts of HT grown emerald are sections under regeneration surfaces R with a thickness of h; the areas between the seed plate and the regeneration surfaces R, excluding the wedgeshaped sections under a and s faces. The total volume of the useful part of the grown crystal thus increases with the seed length, L, and thickness of crystallized layer, h. The length of a seed plate L = 190 mm is in our case limited by the length of the autoclave’s chamber (see Figure 2). Despite the obvious advantage of growing the crystal from a single long seed, it is becoming increasingly difficult to acquire high-quality natural material in such length, while the original seeds are being fragmented to smaller pieces by handling over several cycles of HT growth. As the problem may seem trivial, there are no simple technological solutions. Using crystal fragments, and gradually increasing their length through the HT growth is not feasible, because the growth velocity along the c-axis is close to zero. So, the only solution appears to be the use of multiple seeds aligned in a sequence. Simple geometrical consideration demonstrates that the substitution of a single seed of length L with k shorter seeds of the same cumulative length decreases the proportion of the useful crystal parts: Mk

M1

≈1−

(k − 1) ⋅ [h /tg(15o ) + h /tg(30o )] (k − 1) ⋅ h ⋅ 5.43 ≈1− o o L − h /tg(15 ) − h /tg(30 ) L − h ⋅ 5.43 ACS Paragon Plus Environment

(1) 2

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where M1 is the mass of the useful component in a crystal grown on a seed of length L; Mk represents the total mass of the useful parts in k crystals of the same cumulative seed length; and h is the thickness of a grown layer. This formula demonstrates that under the same conditions (total seed length L = 190 mm and h = 6mm) substituting a single seed plate with two, three or four seeds, this results in a decrease of the useful component to 79%, 59% and 38% compared to the single plate reference sample. To overcome these problems, we present an efficient way for lengthwise splicing of two or more emerald crystals into an almost flawless single crystal, with virtually a full length of useful crystalline material, except for the narrow planar discontinuities at the splicing fronts. Previous experimental studies suggested a possibility of using this method in beryls. Thus, it has been shown that with KDP (KH2PO4) crystals a single growth front can be created between two crystals with the same orientation14. The current paper presents the first successful attempt of splicing two beryl crystals into a bulk single crystal and explores the morphological and nanostructural features in the splicing area. It then discusses the driving forces of this process. Figure 2. Layout of the autoclave system for hydrothermal growth of beryl with the following constituent parts (the indicated dimensions are in millimeters): 1. assembled autoclave; 2, 3. top and bottom thermocouples fixed to the obturator and base, respectively; 4. charge of natural beryl; 5. baffle, dividing the total volume into two zones: growth (top) and dissolving (bottom); 6. frame, fixing two rectangular seed plates; 7. mounting holes drilled into the seed plates.

N.B. : The seeds are fixed with small (1 ÷ 2 mm) overlapping one over another.

Experimental methods Hydrothermal growth experiments Emerald crystals were grown by recrystallization of natural beryl under hydrothermal conditions, with underneath heating and on top cooling of the autoclave. For the mineralizer, we used the multicomponent acidic fluoride- and chromium-containing aqueous solution, commercially used to ACS Paragon Plus Environment

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production of «Russian» emeralds2. The apparatus layout is shown in Figure 2. For the charges of natural beryl we used crystals from Izumrudnye Kopi (Ural, Russia) fractured into 7–10 mm pieces. The seeds plates were of almost rectangular shape and of ~1 mm thickness cut out from previously __ grown emeralds on seed plates with orientation sub-parallel to {5.5.10.6} . Across crystals, the seeds had a standard deviation from the absolute orientation of σϕ ≈ σρ ≈ 1° in polar angles ϕ and ρ. Each of the experimental series used 7 autoclaves placed in shaft furnaces, with 6 autoclaves forming vertices of a regular hexagon around the central (7th) one. The latter served as a baseline for temperature regulation, which was performed using two chromel-alumel thermocouples: one in its base (Tb) and the other – in the obturator (Tu) of the autoclave. Using the PID temperature control system, the temperature was maintained so as to meet the conditions Tav = (Tb + Tu) / 2 = 618±0.1°C, ∆T = Tb - Tu = 75±0.5°C. Inasmuch in all the series half of the autoclave’s free volume was filled with aqueous solution, thus providing equal pressure of ~1.5 kbar under the target temperature Tav = 618°C. The total duration of the growth process across the series varied from 20 to 23 days. Investigations of hydrothermally grown emerald crystals. Resulting crystals were rinsed for 5 minutes in a 45% HF solution to remove a thin amorphous SiO2 film from their surface – without etching the beryl. The spliced crystals were then registered with macrophotography (see Figure 3), and then measured for how much the m faces of crystals deviated from parallel orientation (angle θ in accordance with the definitions on the Figure 3a), using theodolite reflecting goniometer with ~0.25° precision. This degree of error is due to sculpture roughness of m faces generally observed in crystals grown under HT conditions4. The angle ξ between the normals to faces s1 and a2 (see Figure 3a) was measured in some cases also♥. Macrophotographs were recorded on a digital camera, an Helicon Focus 5.1 software was used to increase the focus depth. The morphology of a and s intergrowth faces was explored using optical method with phase contrast on Olimpus BX-51 microscope, and also using scanning electron microscopy (SEM) with Mira-3LMU (TESCAN). Several crystals were then studied for the bulk chemical composition of a layer between the seed and surface R. Major and minor elements analyses were performed by solution ICP-AES and ICPMS, respectively. These were carried out using IRIS Advantage (ThermoJarrell Intertechs Corporation and ELEMENT Finnigan MАТ, Germany) at the Analytical Center for multi-elemental and isotope research of the Sоbоlеv Institute of Geology and Mineralogy, Siberian branch of Russian Academy of Sciences (IGM), Novosibirsk. Instrumental operating conditions and analytical procedures follow those in15. The BHVO-2 standard was used for external calibration. These analyses had an ±1% error for major, and ±5% error for minor elements. Optical microscopy of the internal structure in the splicing area on polished thin (< 0.5 mm) sections of the sample cut parallel to the m-faces was performed using an Olimpus BX-51 microscope. Finally, high-resolution transmission electron microscopy (HRTEM) combined with Fourier transform (FFT) analysis was implemented to determine crystal orientation and nanostructural features of spliced regions. Chemical variations in the intergrowth region were measured _by energy disper_ _ sive spectroscopy (EDS). The interfaces were investigated in two projections, [ 1100 ] and [ 2 4 23 ]. Samples for TEM analysis were cut from the contact region and then mechanically thinned to a ♥

In the case of splicing of the crystals in the perfectly parallel orientation θ = 0°, ξ = 45°. It is easy to check that for a small rotation (first degrees) of one crystal relative to another on polar angles ∆ϕ and ∆ρ (the first is the rotation around the Z-axis, the second is counted from it) the inequality ∆ϕ 2 + ∆ρ 2 < θ 2 + (45° - ξ )2 hold true, with the difference between the left-hand and right-hand sides of the inequality does not exceed 20%. Therefore, further we will use the value θ 2 + (45° - ξ )2 as the upper estimation of the nonparallel intergrown of the two spliced crystals (∆ϕ 2 + ∆ρ 2).


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Figure 3. Intergrowth of emerald crystals with varying angles of misalignment (θ,ξ ). (a) large misorientation (for that crystals θ exceeds 9°) between the seeds results in weak intergrowth of crystals. The morphology of a and s faces in the splicing area is not altered and resembles that on the free end of the crystals. (b-d) two crystals are spliced in quasi-parallel orientation (θ and 45° - ξ are less then 0.25°). Development of vicinal faces on inner a and s faces can be observed as an indication of improving regeneration of the interfacial area. ACS Paragon Plus Environment

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500 µm overall thickness. The specimens were glued onto a copper support and then wedgepolished perpendicular to the contact region using a tripod polisher (Model 590TEM, EMS, Hatfield, PA). Polished thin sections were then finally ion milled using a 2-keV Ar+ ions, and polished at 500-eV (PIPSTM, Gatan). TEM investigations were performed using a 200 kV field-emission TEM microscope (2010F, Jeol) with an ultra-high resolution pole-piece, achieving a resolution limit of 0.11 nm. The microscope is equipped with an EDS spectrometer (ISIS 300, Oxford Instruments, UK) that was used for chemical analysis of intergrowth features.

Results Morphological characteristics of two crystals splicing into a parallel intergrowth During the experiments we observed two distinct types of crystal splicing (Figure 3). The most common case was the irregular intergrowth of two crystals, as shown in Figure 3a; here the angle θ between the normals to faces m of the crystals is ~ 9.5°. These irregularly intergrown crystals are easily separated under a minor withdrawal force, demonstrating a classical feature of nonparallel intergrowth16. Naturally, the elements of growth sculpture on faces a and s, both adjacent to the splicing interface (a1 and s2, in accordance with the definitions on the Figure 3a) and those on other parts of the crystal (a2 and s1), have similar morphology. In this case the average growth rate of the faces a and s do not depend on the localization of these faces with respect to the splicing interface. Significantly less abundant variants of intergrowths are examples depicted in Figures 3b-d (~10% of the total number of intergrowths). The case presented in Figure 3d, in particular, demonstrates the successful and almost complete solution of the applied task formulated in the Introduction. These intergrowths are significantly more durable, requesting an essential load to break the splice, and the crystals do not fracture along the splicing interface. The main difference of these cases from the case in Figure 3a is characterized by the violation of the smooth growth front that is replaced by growth hillocks on the faces adjacent to the intergrowth boundary. Rough topography with large amplitudes and considerable deviations from the expected s and a crystallographic forms are the main characteristics of intergrowths approaching single crystal orientation (Figure 4). Thus, there is a clear change in the growth mechanism from lateral intergrowth4, 9 towards the normal regeneration, characterized by rough R-faces. It is interesting, that the faces of the same crystallographic forms a and s that are simultaneously growing far from the splicing interface (a2 and s1) continue to evolve according the lateral growth mechanism. Morphological features on growing surfaces adjacent to the intergrowth boundary resemble the "torch-like" growth sculptures that expand from twin boundaries on K(H4B5O10)·2H2O crystals17. It is important to note that examples shown in Figures 3b-d and Figure 4 have been observed for almost parallel intergrowths only, for which angle θ is less than 0.25° (below the measurement error). In the cases when we measured the angle ξ for the intergrowths similar to those shown in Figures 3b-d, the value 45° - ξ usually did not exceed 0.25° too. Only in one case the mild complication of the relief on the faces of the a and s around the area of splicing was observed for the crystals intergrowth with a misalignment of ~ 0.4° So, we can conclude that vicinal complications on the faces a1 and s2 adjacent to the border of splicing are observed for the parallel splices of crystals only. The opposite statement is true too. One can ask, what is the reason of the evident morphological differences of the surfaces near the intergrowth interface for almost parallel splices, shown in Figures 3b – 3d? We see different variants of the answer to this question. For example, it is possible to explain by the different amplitude of deviations from intergrowths parallelism, that we fail to fix by our goniometry methods. Or we can assume that the events of splicing for different samples occurred at different point in time of growth process: for the sample in Figure 3d the splicing began at the start of the growth process,


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Figure 4. Rough relief of splicing interface from Figure 3c, produced by large amplitude vicinals on adjacent s (right) and a (left ) faces. The blue vertical line indicates an approximate position of turbulent splicing interface (macrophotography). The inset shows an SEM image of splicing interface.

whereas for the sample 3b it started at the very end, and advanced regeneration surface just did not have time to form. The final answer to this question is the task for future research. Detailed analysis of the splicing region and investigation of the local singularities and chemical composition of this region was further studied on perfectly intergrown samples, such as shown in Figure 3d. The thickness of grown beryl h on each side of the seed is the same, and amounts to 6.5 mm. The total duration of the growth process is 23 days, and hence, the average growth rate of the R surface is 0.28 mm/day. To conduct local studies of the composition of the crystal near the intergrowth boundary by EDS method should be considered that the quantitative chemical analysis of beryl by this method is not possible, since it does not define light elements Be, Li and water. Therefore, as a base for EDS measurements, we used the chemical analysis of the part of grown layer away from the region of splicing by the "wet" chemistry methods, averaged over the crystal's thickness. Bulk chemical composition of grown beryl measured by ICP method is (in wt.%): SiO2 – 64.88, Al2O3 – 15.79, BeO – 13.68, Fe2O3 – 3.09, Cr2O3 – 0.65, NiO – 0.19, MgO – 0.04, MnO – 0.04, CuO – 0.02, Li2O – 0.05, Na2O – 0.04, K2O – 0.09, LOI (loss on ignition) – 1.30, Total – 99.87. The source of trace elements (Ni, Mn, Cu and part of Fe) is the autoclave's steel. Recalculation of the chemical composition on the base of 18 oxygen atoms, neglecting the low concentrations of Cu and Mn, and assuming that all the loss on ignition is due to loss of water, gives the following bulk composition for our emerald crystals: Be3.01(Al1.72Fe0.21Cr0.05Ni0.01Mg0.01)Σ=2(Si5.98Be0.02)Σ=6O18·(Li0.02Na0.01K0.01)Σ=0.04(H2O)0.40 (2) Intergrowth structure of the splicing area. The splice similar to that shown in Figure 3d was used as the object for studies of the internal structure of the intergrowth area. Its bulk chemical composition is featured by (2). Figure_ 5 shows a micrograph of the splicing area of polished thin section of the sample, cut in parallel to (1100) . The photo clearly presents the projection of the splicing interface (defined in Figure 5 as Interface-I and _ Interface-II) on the plane (1100) , the output of which on the growth surface is __ completed by the reentrant angle, resembling that of twin intergrowths. In the projection on (5.5.10 . 6) the interface between two crystals looks as serpentine-like line (Figure 4). The cracks near the splicing interface ACS Paragon Plus Environment

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were produced during the preparation of the thin section, while releasing tensions, concentrated in this area. The intergrowth region between spliced beryl crystals comprises different growth sectors __ that are formed in three distinct growth regimes: (i)_ growth on regeneration surfaces R {5.5.10.6} , _ (ii) intermediate growth stage above the inner a {1120} and s {1122} faces, and (iii) rapid growth of the central part __that __ closes up the intercrystal gap, producing subvertical splicing interfaces roughly following {3.3. 6 .10} planes. These sectors in Figure 5 are marked by the capital letter G, the subscript symbol in which indicates the growing surface that forms the sector; e.g. Ga2 is the growth sector above a plane of Seed-2. The subscript symbol f in the notations Gf1 and Gf2, are an abbreviation of the word “fill”. An interface between two different growth sectors, for example, between Ga2 and GR2, is denoted as Ga2 | GR2. In an ideal case, when the seeds are perfectly aligned, growth sectors below the seeds are related to the growth sectors above the seed plane by two-fold axis, represented in the group of symmetry P6/mcc. From Figure 5 this qualitatively appears to be true, with some morphological differences. For example, the unique presence of Gs2 in the upper section is explained by the fact that it wasn’t formed during this bicrystal growth experiment; it was inherited as a part of the crystal from which Seed-1 was prepared. The angles between individual linear structures (i.e. boundaries of the seeds, projections of the faces a and s, the boundaries between the growth sectors), measured from Figure 5, are listed in Table 1.

Figure 5. Thin section of the splicing area of fully intergrown beryl crystals (see Figure 3d); imaged with partially crossed Nicols. The emerald crystal (above the thin_ section; dark line intersecting the crystals are seeds used for HT growth of emerald) was cut along the m (1100) plane. The field of _ view, displaying the intergrowth structure, is outlined. Denotations a1, s2 and R1 and R2 indicate inner (11 of_ Seed_ __ 20) prism _ face _ 1, inner (1122) bipyramidal face of Seed-2, and both regeneration surfaces (5.5.10.6) and (5.5.10.6) , respectively. Letter 'G' denotes growth sectors above specific crystallographic faces related to the particular seed crystal; e.g. Ga2 is the growth sector above a plane of Seed-2. The splicing interfaces are produced between Gf1 | Gf2 growth sectors. Different interference colors on the Figure are explained by the variation of the local refractive indices caused by tensions.


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Table 1. Angles (∠) between the linear structures delimiting the growth sectors, measured from Figure 5. Only the interfaces in the upper section are listed, the situation in the lower section is related by two-fold axis symmetry. Row # 1 2 3 4 5 6 7 8 9 10

Linear structure (see notations in Fig. 5) Regeneration surface of Seed-1 (R1) Regeneration surface of Seed-2 (R2) Original inner a-face of Seed-1 (a1) Original inner s-face of Seed-2 (s2) Lower interface between GR2 and Ga2 Upper interface between GR2 and Ga2 The boundary between sectors Ga2 and Gf2 The upper splicing interface (Interface-I) The boundary between sectors Gs1 and Gf1 The boundary between sectors GR1 and Gs1

∠ to R1 0.0° 0.9° 29.2° -15.9° 29.4° 46.2° 40.2° 86.4° 127.4° 150.4°

Near-perfect alignment of the seed plates, embedded in HT grown product (see the final emerald bicrystal in the upper part of the optical micrograph in Figure 5), can be estimated by measuring the angle between the seeds, which corresponds to ~ 0.9° (row#2 in Table 1), suggesting a small misalignment between the two crystals. However, surprisingly, the angle ξ measured between the outer a1 and s2 faces equals to 45.1° (difference between the rows#3 and #4 in Table 1), which closely corresponds to the theoretical value (45.0°) between these two faces for single beryl crystal, indicating that the two crystals are in fact intergrown in identical crystallographic orientation. Within the measurement error of the angles, the crystals appear to produce a perfectly parallel splice7. Now let's consider the structure of the various growth sectors. Growth sectors above the regeneration surfaces R (marked as GR1 and GR2 in Figure 5) are visible as a subtle zig-zag horizontal stripes (at higher magnifications), are result of spontaneous growth zonality from the pulsatile nature of the HT growth process18. Each band indicates one time-lapse during the crystal growth. Optical heterogeneities that extend in subnormal direction across these bands are micro-facets of the ridges on the regeneration surface R 4, 5, 11, 18. Fourier transform analysis via short-time sliding along the profiles perpendicular to R-faces showed that, at least, the first 2 mm of beryl grown in GR sectors are characterized by constant fluctuations in optical density (~15 fluctuations per 1 mm), indicating growth mode of a constant rate. Then let us consider growth sector Ga2 and the boundary Ga2 | GR2. Growth zonality in the sector Ga2 is quasiparallel; the latter suggests that all this sector is really formed by the same growing face a. It is known, that the slope of the boundaries between growth sectors above these two faces is uniquely determined by the angle between the normals to these faces, and the ratio of their growth velocity19. In our case, this relationship is determined by the equation: tg( β ) =

Va sin(α ) , VR − Va cos(α )

(3)

where α is the angle between the normals to faces R and a, and β is the angle between the normals to the face a and their interface. In our case, the interface Ga2 | GR2 consists of two line segments with the obvious kink between them. From the above equation and the data listed in Table 1 one can find that at the breaking point the ratio between the growth rates VR : Va changes almost abruptly from ~7:1 [see 13] up to ~7:3, while the rate of the regeneration surface R remained constant. On the basis of the average value of VR = 0.28 mm/day this change in growth regime occurs after about 5.5 days of HT growth. Earlier, we noted that parts of the bicrystal below and above the seeds should be related by two-fold axis. However, despite the similarity of the sectors Ga1 and Ga2, the latter exhibit some quantitative differences. Accordingly the calculations that is similar to the ACS Paragon Plus Environment

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above, kink of the interface Ga1 | GR1 (and, therefore, the change of the growth regime of the face a) dates to ~ 2 days since the beginning of the growth. The reasons of these quantitative differences between Ga1 | GR1 and Ga2 | GR2 _may be due to different initial shape of the seeds – the right end of Seed-1 is formed by face a {112 0} but the left end of Seed-2 is cut by the face c {0001} , growing with almost zero speed. Growth sectors Gf1 and Gf2 that produce a splicing interface Gf1 | Gf2 from the adjacent seed and close up the intercrystal gap also contain some growth features. Rather interesting is the fact that some growth features run across both growth sectors. That fact means that sectors Gf1 and Gf2 are actually common growth structure. This is consistent with Figure 3d that presents the similar splice – large subindividuals over the intergrowth boundary form a common growth front without be broken down to the fragments on either side of the splicing interface. Interweaving of protrusions from the opposite growth sectors may explain the large load that is necessary for such spliced interface to be broken apart. Note that the Gf1 | Gf2 interface (Interface-I) from upper section is displaced with respect to its lower counterpart (Interface-II) for the amount of overlapping seed displacement. Both splicing interfaces are sub-vertical, with a small inclination (approx. 3°) towards the parent seed (see Figure 5). It is also a surprising fact that the interface Ga2 | Gf2 (and Ga1 | Gf1) is almost a straight line, since the growth velocity of the a-face was low during the first 5.5 days, and then its rate increased by the factor of 3, while the rate of the regeneration surface R remained constant. The explanation for this can be also non-constant growth rate of the surface forming the growth sector Gf2 (Gf1), changing in proportion to the growth rate of the face a. From the slope of Ga2 | Gf2 (Table 1) we may conclude that the growth rate of the surface forming the growth sector Gf2 approaches the growth velocity of the surface R through 5.5 days. _

Let us consider now the group of the growth sectors of the faces s {112 2} . These faces were not present in the original surface of the seeds near the region of splicing, unless one doesn't take into account the "trace" of s that is the part of Seed-2 (Gs2). However, there is the growth sector in the overgrown layer that corresponds to face s, i.e., Gs1. The growth zonality in this sector is not a straight line, as in the discussed above GR2, Ga2, and GR1, but the curves that is approaching to the orientation of the face s2 near the Interface-I and become more flat far from it. From the inclination of GR1 | Gs1 (Table 1, Row#10) according to equation (3) one can find the ratio VR : Vs1 ≈ 2 : 1. This _ value is fundamentally different from that for the faces s {112 2} far from the splicing boarder (7 : 1). This demonstrates the accelerating effect of the Interface-I on the growth of the face s1. It is interesting, that the sector Gs2 in position under the Seed-2, which is symmetric to the sector Gs1 located above the Seed-1, is absent. This, apparently, is explained by the asymmetry of the plot of contact of Seed-1 and Seed-2.

Fine structure of the splicing interface. To understand the mechanism of crystal splicing, the intergrowth area (Gf1 | Gf2 interface) of the sample shown in Fig. 3d was studied under TEM. In the splicing area we observe a high density of long pores, running sub-parallel to the interface from both domains (Figure 6a). Pores are often parallel and have a constant width of 12 ±2 nm (Figure 6b). Given the fact that the density of pores in the splicing area increases quite suddenly, they must be directly related to splicing of the two oppositely growing domains. An important indication of their nature is shown by chemical analysis of their composition. As the Si/Al ratio for bulk beryl by EDS method is identical to the ratio measured by ICP (Eq. 2) we can take this as a reference for calculating compositions of bulk beryl, interfacial beryl and amorphous material in pockets, as listed in Table 2.


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Figure 6. TEM study of splicing interface in projection normal to m-face, showing (a) interwoven bands of beryl domains separated by multiple faults running sub-parallel to the Gf1 | Gf2 interface. (b) Channels of amorphous material intersect crystalline beryl. The inset shows characteristic dark bands along the axial section of the channels; some channels extend along the viewing direction (marked by arrows) (c) HRTEM image of channels with a characteristic dark core due to local enrichment in Si content in the upper-right part of the HRTEM image. In beryl, no misfit can be observed_ across the domains. (d) FFT image of crystalline beryl confirming [ 1100 ] projection.

Chemical composition of pore-free bulk beryl is identical to the beryl from HT grown sections above R-surfaces with Si/Al elemental ratio of 3.47 (see Eq. 2). However, when we analyze crystalline areas between the pores, the composition of beryl becomes Si-rich, with Si/Al ratio increasing to 4.42 (see Table 2). However, this radical shift in compositions not accompanied by significant change in crystal structure, which remains beryl-like. Such violation of stoichiometry is permissible, if we consider that the main components of beryl Si, Al and Be allow limited substitution within _beryl structure20. We can see that some pores are also running parallel to viewing direction [ 1100 ] (Figure 6c), suggesting that in the splicing area the pores adopt different directions. They are hosted by fault-free crystalline beryl that except for chemical variations shows no change in crystallographic orientation, nor it contains planar defects or misfit dislocations that would indicate an intergrowth.


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Table 2. EDS analysis of various parts of beryl near splicing interface expressed (at. %).

Bulk beryl* Interfacial beryl* Amorphous mat.

Si

Al

Fe

Cr

Ni

O

Si/Al

26.08 ±0.16 27.40 ±0.20 29.82

7.51 ±0.20 6.20 ±0.23 3.70

0.92 ±0.04 0.70 ±0.03 0.41

0.22 ±0.01 0.18 ±0.01 0.09

0.07 ±0.01 0.04 ±0.02 0.03

65.20 ±0.03 65.47 ±0.05 65.96

3.47 4.42 8.06

* Values with std. deviation include 3 measurements.

_

_

Figure 7. TEM study of splicing interface in [ 2 4 23 ] projection. (a) Bright-filed TEM image showing pockets of amorphous material lined along the interface between splicing crystal domains. Microphotograph of emerald bicrystal in the upper part of the image shows an approximate viewing direction. _ (b) HRTEM image of the area indicated by while dotted square in Figure 7a shows well resolved { 10 1 0} and __ {0112} planes with lattice spacings of 7.98 and 3.98 Å. Across the amorphous pocket lattice planes are _ discontinued and shifted for ½ (1010) . Indicated stripes of HRTEM image demonstrating tilt and shift of _ (1010) lattice planes are shown below in images (c-e). (f) Fourier transform image (FFT) of (b) showing _ _ splitting of reflections, particularly those with small d-values, i.e. {22 4 4} (outlined). In the left part of the FFT image, simulated electron diffraction patterns calculated for the tilt of 0.4° are superimposed (red – Domain I, green – Domain II). The rings in the background are convolution of the amorphous material with the contrast transfer function of given imaging conditions. (g) Inverse FFT image of the upper sec_ tion of HRTEM image shown in (b), using {1010} reflections. ACS Paragon Plus Environment

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Even more interesting is the composition of amorphous material within the channels protruding the interfacial beryl. EDS analysis of amorphous material measured near the edge of Si-rich beryl showed a further increase in Si content, with the Si/Al ratio raised above 8. The increase in Si is accompanied by a decrease of other constituents (see Table 2). Such a high deviation in stoichiometry breaks down the crystalline structure and the material becomes amorphous. Another interesting features are dark, ~3 nm wide bands extending along the center of the channels of amorphous material. Their composition indicates pure silica that is the remnant of HT solutions. If we anticipate that the excess of O atoms is compensated by the corresponding number of Be atoms (undetectable by EDS) we obtain approximate composition of the amorphous phase is BeO⋅Al2O3⋅16SiO2. In other words, the amorphous phase in the pores approaches the composition of pure silica gel. _

_

To verify for misalignment component (φ) the interface was additionally investigated in [ 2 4 23 ] projection. The position of the interface was identified by the line of voids enclosed by the splicing domains, as shown in TEM image in Figure 7a. The voids lined along the interface are occurring at 500-700 nm, suggesting that they are result of accumulated misfit between the two crystal domains. A close view of the interface shows_ that the voids are filled with amorphous material, such as those in Figure 6. In this projection, (1010) lattice planes crossing the interface are tilted from 0.0-0.8. The largest tilt is measured in the upper part of the image, above the amorphous pocket, gradually _ progressing into continuous lattice with no tilt of (1010) lattice planes. Typically, the areas of worst tilt contain an amorphous pocket, whereas regions between two pockets show almost no tilt. This _ change in tilt can be best seen if we glance along the (1010) planes in different sections of the interface. Above the pocket, the lattice planes are abruptly tilted by 0.8° when crossing the interface (Figure 7c). Soon after this _ maximum tilt amorphous pocket occurs where accumulated tilt is re7d). After the pocket is laxed, including the ½ (1010) shift, while the tilt is reduced to 0.4° (Figure _ closed, the lattice is continued with an almost perfect translation of (1010) lattice planes from one into another domain (Figure 7e). Fourier transform (FFT, Figure 7f) of lattice image across the interface shows two sets of reflections that are rotated _ by _ ~0.4°. Splitting is best observed on more distant reflections with smaller d-values, such as {224 4} . FFT image can be used to select specific reflections to visualize lattice translations, such as strain fields and phase shifts across interfaces21. _ In our case we use {1010} lattice planes. The resulting inverse Fourier image (Figure 7g) shows _ both, large tilt in the upper part of HRTEM image and ½ (1010) shift across the amorphous pocket.

Discussion In the following we discuss the possible mechanism of crystal alignment and intergrowth, and the origin of amorphous pockets that form along the splicing interface.

Self-alignment of crystals during the formation of parallel intergrowth. Bicrystal that was analyzed in our study (Figure 3d, Figure 5), had Seed-1 and Seed-2 misaligned for 0.9°. Taking into account that during the growth process the seed plates must be positioned vertically (Figure 2), the deviation of the seed of length L from the vertical around the top mounting point for an angle 45 − ξ = ± 0.9° requires work A = (1 – cos (45 − ξ))⋅p⋅L/2, where p is the weight of the seed in the medium with a specific gravity ~0.6 g/sm3. For the seed with the dimensions of 8×2×0.1 cm at a specific gravity of beryl 2.7 g/cm3 A ≈ 1.3⋅10-5 J. Of course, the above value is just an estimate, because during the hydrothermal process in the working chamber of the autoclave there are acting turbulent flows that complicate the situation18,22, by rotating the seed from the vertical position. On the other hand, the contribution of these turbulences can also be productive, reducing the angle of misalignment. Thus, we can suggest that there are certain driving


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forces that are committing a substantial work aimed to agitation of the seeds around the crystallographically parallel orientation. One can come to the existence of such driving forces from the standpoint of statistical consideration of sub-parallel intergrowths (within 0.25°) among the total recorded number of spliced crystals. Let us define the probability of an event that misorientation of two seeds (expressed in the values of polar angles ρi and ϕi) will not exceed some certain value a♥. It is necessary to take into consideration that the standard deviation of misalignment σϕ ≈ σρ ≈ σ = 1° (see Experimental section) for both seeds, and the mathematical expectations E[ρ1]= E[ρ2], E[ϕ1]= E[ϕ2] The probability of such event can be defined as:

  ρ − ρ 2  2  ϕ1 − ϕ 2  2 a 2  P ( ρ1 − ρ 2 ) 2 + (ϕ1 − ϕ 2 ) 2 ≤ a 2 = P  1 + ≤   2σ   2σ  2σ 2  

(

)

(4).

It can be shown that if the random independent values ρ1, ρ2, ϕ1, ϕ2 have the same normal distri2

2

 ρ − ρ 2   ϕ1 − ϕ 2  2 bution, then  1  +  ∈ χ (2) is chi-square distribution with two degrees of freedom.  2σ   2σ  Thus, the required probability P ( ρ1 − ρ 2 ) 2 + (ϕ1 − ϕ 2 ) 2 ≤ a 2 can be found from the χ2(2) distribution. The probability values calculated for different values of a are given in Table 3.

(

)

The comparison of experimentally determined frequencies of formation of parallel intergrowths among the total recorded number of intergrowths of crystals to ~ 10% with those calculated (Table 3) for the case a = ¼σ = 0.25°, which is 1.6%, shows the difference in almost one order of magnitude. In our calculations we initially assume independence of orientations of the two seeds, i.e., the absence of some driving forces, seeking to "correct" the mutual orientation of the Seed-1 and Seed2 to a parallel alignment. The difference of one order of magnitude between the experimentally determined and the calculated probabilities shows the contradiction with our original assumption and suggests the presence of driving force, that significantly improves the initial alignment of the seeds. Note that actually, instead of value ∆ϕ 2 + ∆ρ 2 we use their upper estimations θ 2 + (45° - ξ )2. It means that the real differences between the experimentally determined frequencies and calculated values will be even higher. Table 3. The probability of the event that the differences in orientation of two randomly selected seeds (expressed in the values of polar angles ρ and ϕ) will not exceed a certain a. a

σ = 1°

½σ = 0.5°

¼σ = 0.25°

⅛σ = 0.125°

¹⁄16σ = 0.063°

P(ρi,φi)

0.221

0.0606

0.0155

0.0039

0.00098

We failed to find in the literature an evidence of self-alignment of macroscopic crystals, however this phenomenon is intensely studied in nanocrystals23. Studying self-assembly and growth of rutiletype TiO2 mesocrystals from Ti-butoxide colloidal solutions24, demonstrated that rutile fibers selfalign due to yet unknown electromagnetic interactions that convey the essential structural information between the nanocrystals. Similarly as in our study, the success of self-alignment is enhanced by thermal agitation of crystals. Given the fact that our seed crystals are loosely attached in position, the crystals have a chance to experience similar electromagnetic interactions as described for nanocrystals, when they approach a critical distance during growth. If so, this could explain the ♥

From a mathematical point of view, this task is identical to the test of the well known statement »The artillery shell doesn't fall in the same crater twice«.


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observed near perfect self-alignment of HT beryl crystals when under certain statistical angular misfit (example shown in Figure 3d). If the initial misalignment is larger, the work needed to tilt the crystals into crystallographically aligned position is too large and the crystals stay misaligned (examples in Figures 3a-c). The reason, why above certain tilt of macroscopic crystals can not be selfaligned, is their large mass that involves an increasing work load to align them. So, efeectively, there exists a critical tilt for every given size of crystals that can be overcome by short-range interaction forces under given HT growth conditions.

Intergrowth of beryl crystals by splicing. The main macroscopic characteristic of near-perfect and poorly aligned bicrystals is overgrowth on inner a- and s-faces. In poorly aligned bicrystals, there is almost no overgrowth on these faces, whereas in well aligned crystals it is abundant. In the first case (shown in Figure 3a) the adjoined crystals have well developed, smooth, inner a- and s-faces that produce a re-entrant angle of 135°, similar to that of twinned crystals. The growth of such a bicrystal is somewhat analogous to the twin-plane reentrant (TPRE) mechanism described for twin boundaries25,26. Certainly, a parallel splice of two crystals cannot be called a twin from the positions of point symmetry because it has unique symmetry elements that are different from those of a twin27. However, if we consider the twins from the standpoint of space symmetry, we may add simple translational elements, such as shift of 1/2 of the unit cell and obtain the same result28,29. In this case, however, the differences between sub-parallel splice and a twin are not that significant. From this point of view a sub-parallel splice and a twin in its mature state of growth30 should develop similar re-entrant angles. Completely different situation is developed when crystals are met near identical crystallographic orientation (Figures 3b-d). Better initial alignment of beryl crystals evidently leads to a sharp increase in relative growth rates of faces a and s adjacent to the interface from 10÷15 to 30÷90, approaching that of the regeneration surface R (VR = 100), while the outer a- and s-faces remain in the range of 10÷15. Overgrowth on inner (re-entrant) a- and s-faces is a result of rapidly developing 1st generation (Gs1 vs. Ga2) and 2nd generation (Gf1 vs. Gf2) splicing sectors, that eventually close-up the intercrystal (re-entrant) gap. It remains unclear why near identical crystallographic orientation regeneration of the re-entrant faces occurs at such a high rate, which, from the point when the two crystals are met in this orientation, is accelerated to match the fastest growing regeneration surface R. This especially holds for Gf growth sectors that produce the splicing interface. This demonstrates how splicing of two seeds near identical crystallographic orientation, accelerates the growth of otherwise slowly growing directions normal to a- and s-faces even beyond the growth rate of the joint regeneration surface R at the onset of Gf sector formation. Moreover, rapid regeneration of reentrants is related by two-fold axis presented in the group symmetry of beryl P6/mcc and the growth conditions is thus the same for both sides. The detailed consideration of this issue is the task of future research.

Formation of amorphous pockets along the splicing interface. At this stage, it is difficult to give a definite answer about the formation of amorphous pockets in the splicing area, however, based on TEM results we believe that the reason for their occurrence is two-fold: (i) stoichiometry breakdown that determines their composition and (ii) lattice strain-filed due to misalignment of splicing domains controlling their spatial distribution. Growth of crystals from a quasi-liquid phase with substoichiometric solutions has been described by the VLS mechanism31. If the composition of the quasi-liquid phase is close to the stoichiometry of beryl, it will continue to crystallize as normal beryl, however when the stoichiometry of the quasi-liquid phase is not completely identical to the crystal composition, a drift of the solid phase composition could take place at the end of crystallization stage. This process can explain the formation of beryl with changed stoichiometry, such as observed in our case in splicing region. This, however, should lead ACS Paragon Plus Environment

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to local changes of unit cell parameters, which in turn can either enhance or weaken the tensions near the splicing area. Remains of the quasi-liquid phase depleted in major components (Be, Al) are then entrapped by the growth fronts of the crystal in form of amorphous phase pockets. If our understanding of this process is correct, we can expect to detect similar pockets of amorphous phase also in the body of bulk beryl in growth sectors of the growing micro-faces that form the front of Rsurface. Their distribution, however, appears to be related to lattice strain. Namely, the pockets of amorphous residue are more likely to form in the areas of largest misfit between the two crystal domains.

Summary In the present study we demonstrated that splicing of two crystals near identical crystallographic orientation is possible. The study was demonstrated on the example of beryl, Be3Al2Si6O18, growing under hydrothermal conditions. We found that self-alignment of crystals takes place only when the seeds are loosely attached near identical crystallographic orientation within the critical angle of misalignment. This allows the crystals to seek for fine crystallographic orientation as the crystal growth fronts approach the nanometer range. One of the most reliable markers of successful alignment of crystals is complete regeneration of the inner (re-entrant) a- and s-faces on fully grown bicrystals. To our knowledge this is the first experimental evidence of self-alignment of macroscopic crystals that is driven by short-range molecular interactions. The result has undoubted practical importance because it allows to increase the yield of useful product within a single HT growth cycle.

Acknowledgements The research was supported by the Slovenian – Russian bilateral projects BI-RU/14-15-0025 and BI-RU/16-18-004. Hydrothermal experiments were carried out using equipment provided by TAIRUS Company (Novosibirsk, Russia). We thank our colegues Dr. S.Z. Smirnov, Dr. Ju.M. Borzdov, Dr. N.S. Karmanov and O.A. Kozmenko (Sobolev Institute of Geology and Mineralogy SB RAS, Novosibirsk, Russia) for the assistance in the investigations and scientific discussions. We are very grateful to Dr. E. V. Thomas (Higher School of Economics, Moscow, Russia) for the great work on the design of the manuscript.

References (1) Flanigen E.M., Mumbach N.R. Hydrothermal process for growing crystals having the structure of Beryll in an acid halide medium. Pat. USA №3.567.643, 1971. (2) Lebedev A.S., Il’in A.G., Klyakhin V.A. Hydrothermally grown beryls of gem quality (in Russian). In Morphology and Phase Equilibria of Minerals. Proceedings of the 13th General Meeting of the International Mineralogical Association, Varna (Sofia, Bulgaria,1982), 1986, Vol. 2, pp. 403–411 (in Russian). (3) Kiefert L.,Schme tzer K. The microscopic determination of structural properties for the characterization of optical uniaxial natural and synthetic gemstones. Part 2: Examples for the applicability of structural features for the distinction of natural emerald from flux-grown and hydrothermally-grown synthetic emerald. Journal of Gemmology, 1991, Vol. 22, No.7, pp. 427-438. (4) Lebedev A.S., Askhabov A.M. Regeneration of beryl crystals. Zapiski Vsesoyuznogo Mineralogicheskogo Obshchestva, 1984, v. 113, No. 5, pp. 618-628 (in Russian). (5) Schmetzer K. Characterization of Russian hydrothermally-grown synthetic emeralds. Journal of Gemmology, 1988, Vol. 21, No. 3, pp. 145-164.


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(6) Schmetzer K., Schwarz D., Bernhardt H.-J., Häger T. A new type of Tairus hydrothermally-grown synthetic emeralds, colored by vanadium and copper. Journal of Gemmology, 2006, v. 30, No. 1/3, pp. 59-74. (7) Goldschmidt V. Atlas der Krystallformen; Carl-Winters-Universitaetsbuchhandlung: Heidelberg, 1913; Vol. 1, p 182. (8) Gavryushkin P.N., Thomas V.G. Growth Kinematics of the Regeneration Surfaces of Crystals. Crystallography Reports, 2009, Vol. 54, No. 2, pp. 334–341. (9) Demianets L.N., Ivanov-Schitz A.K. The growth mechanism and morphology of hydrothermally grown oxide compaunds: fractal approach. J. Phys.: Condens. Matter, 2004, Vol. 16, p. 1313. (10) Bekker T. B., Barz R.-U. Study of Growth Faces in Hydrothermally Obtained Beryl Single Crystals Using (556)-Orientated Seeds. Crystal growth & design, 2007, 7(9), p. 1898 - 1903. (11) Thomas V. G. Gavryushkin P. N., Fursenko D. A. 2D Modeling of the Regeneration Surface Growth on Crystals. Crystallography Reports, 2012, Vol. 57, No. 6, pp. 962–974. (12) Thomas V. G., Gavryushkin P. N., Fursenko D. A. 2D Modeling of Regeneration Surface Growth on a Single-Crystal Sphere. Crystallography Reports, 2014, Vol. 60, No. 4, pp. 583–593. (13) Thomas V.G., Demin S.P. Regeneration of nonsingular surfaces of beryl as the simultaneous growth of positive and negative crystals. Abstr. Intern. 1-st Conf. «Crystallogenesis and Mineralogy», 2001, p.397398 (in Russian). (14) Zaitseva N., Smolsky I., Carman L. Growth phenomena in the surface layer and step generation from the crystal edges. J. Crystal Growth, 2001, Vol. 222, p. 249–262. (15) Shatsky V.S., Sitnikova E.S., Koz'menko O.A., Palessky S.V., Nikolaeva I.V., Zayachkovsky A.A. Behavior of incompatible elements during ultrahigh-pressure metamorphism (by the example of rocks of the Kokchetav massif). Russian Geology and Geophysics, 2006, Vol.47, p.482-496. (16) Fersman A.Ye. The elements of surface between two simultaneously crystallizing substances // DAN USSR, 1922, A, 7-8 (in Russian). (17) Wojciechowski V.N., Nikolaeva V.N., Velichko I.A. On the specific features of potassium pentaborate crystal growth. Crystallography, 1982, Vol. 27, p.975 – 980 (in Russian). (18) Thomas V.G., Demin S.P., Fursenko D. A., Bekker T. B. Pulsation processes at hydrothermal crystal growth (beryl as example). J. Crystal Growth, 1999, Vol. 206, p. 203–214. (19) Prywer J. Theoretical analysis of changes in habit of growing crystals in response to growth rates of individual faces. J. Crystal Growth, 1999, Vol. 197, p. 271–285. (20) Thomas V.G., Klyakhin V. A. Specific features of incorporation of chromium in the beryl structure under hydrothermal conditions (experimental data). in Mineral Forming in Endogenic Processes, edited by Sobolev, N. V. (Nauka, Novosibirsk) , 1987, pp. 60–67 (in Russian). (21) Hÿtch M.J., Snoeck E., Kilaas R. Quantitative measurement of displacement and strain fields from HREM micrographs. Ultramicroscopy, 1998, Vol. 74, p. 131-146. (22) Thomas V.G., Bekker T. B. The method of testing by temperature fluctuations (TTF) to investigate the heat-mass transfer in autoclave during the hydrothermal crystal growth . Proceeding IV Int. Conf. 'Single Crystal Growth and Heat & Mass Transfer', Ed. Ginkin V.P., Obninsk, 2001, Vol.3, p. 764-772. (23) Yoreo de J.J., Gilbert P.U.P.A., Sommerdijk N.A.J. M., Penn R.L., Whitelam S., Joester D., Zhang H., Rimer J.D., Navrotsky A., Banfield J.F., Wallace A.F., Michel F.M., Meldrum F.C., Cölfen H., Dove P.M. Crystallization by particle attachment in synthetic, biogenic, and geologic environments. Science, 2015, Vol. 349, Issue 6247, p. aaa6760-9, DOI: 10.1126/science.aaa6760. (24) Jordan V, Javornik U, Plavec, J, Podgornik A, Rečnik A. Self-assembly of multilevel branched rutiletype TiO2 structures via oriented lateral and twin attachment. Scientific Reports, 2016, Vol. 6, 24216. (25) Becke F. Uber die Ausbildung der Zwillingskristalle. Fortschr. Mineral., 1911, 1. (26) Tiller W.A. The science of crystallization: microscopic interfacial phenomena. Cambridge University Press, 1995, 392 pp. (27) Putnis A. Introduction to Mineral Science. Cambridge University Press, 1992, 457 pp. (28) Frank F.C. Crystal dislocations-Elementary concepts & definitions. Phylosophical Magazine, 1951, Vol. 42, p. 809 – 819. (29) Vainshtein B.K., Fridkin V.M., Indenbom V.L. Modern Crystallography-II. Spr.-Ver. Berlin – Heidelberg New York, 1982, 436 pp. (30) Drev S., Rečnik A., Daneu N. Twinning and epitaxial growth of taaffeite-type modulated structures in BeO-doped MgAl2O4. CrystEngComm, 2015, Vol. 15, p. 2640-2647. ACS Paragon Plus Environment

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(31) Wagner R. S., Ellis W. C. Vapor-liquid-solid mechanism of single crystal growth. Appl. Phys. Lett., 1964, Vol.4, p. 89-90.


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For Table of Contents Use Only

Crystallographic assembly of macroscopic crystals by sub-parallel splicing of multiple seeds. Thomas V.G., Daneu N., Rečnik A., Fursenko D.A., Demin S.P, Belinsky S.P., Gavryushkin P.N.

We investigate poorly studied phenomenon of self-orientation of macroscopic crystals that has been observed under hydrothermal growth conditions with intention to understand and exploit this mechanism for the formation of parallel intergrowths for industrial production of gem quality emeralds using multiple seeds.


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Crystallographic Assembly of Macroscopic Crystals ... - ACS Publications

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