- Email: [email protected]
Evaluation of improvement strategies of grain structure properties in high performance multi-crystalline silicon ingots
Journal of Crystal Growth 514 (2019) 114–123
Contents lists available at ScienceDirect
Journal of Crystal Growth journal homepage: www.elsevier.com/locate/jcrysgro
Evaluation of improvement strategies of grain structure properties in high performance multi-crystalline silicon ingots
T
⁎
M. Trempaa, , C. Kranerta,b, I. Kupkaa,b, C. Reimanna,b, J. Friedricha,b a b
Fraunhofer IISB, Schottkystr. 10, 91058 Erlangen, Germany Fraunhofer THM, Am St.-Niclas-Schacht 13, 09599 Freiberg, Germany
A R T I C LE I N FO
A B S T R A C T
Communicated by A.G. Ostrogorsky
High performance multi-crystalline silicon (HPM-Si) for the use in photovoltaics is characterized by a very fine grain structure and a high content of random grain boundaries, finally resulting in a low dislocation density and consequently in a high material quality. Typically, the grain size increases and the fraction of random grain boundaries decreases over ingot height due to annihilation mechanisms, especially in the first 150 mm. One approach for further material improvement is to further increase the initial random grain boundary fraction and to maintain it as high as possible over the complete ingot height. In this work, several theoretical approaches to achieve these points were evaluated by experiments in G1 scale. Firstly, the influence of the silicon seeding material on the initial grain structure was investigated regarding the effect of extremely fine Si particles in the µm to nm range and the bulk density of the particle layer. Secondly, the effect of the initial geometrical grain boundary arrangement in the seed layer was evaluated. For that purpose, special seed alignments similar to the Quasimono approach were tested. Finally, the process parameter growth rate was varied in a wide range to investigate its influence on the evolution of the grain boundaries during growth. The results show that the optimum for the initial random grain boundary fraction is already reached by the existing methods/commonly used seed materials. Concerning the decrease of the random grain boundary fraction over ingot height, some technical aspects were identified which are able to keep the amount of random grain boundaries at a high level. However, the practical realization within an industrial setup might be difficult.
Keywords: A1. Directional solidification A1. High performance multi-crystalline silicon A1. Grain structure A1. Grain boundaries A1. Recombination active defects
1. Introduction Multi-crystalline (mc) silicon produced by directional solidification technique is still the most used material for the production of silicon solar cells with a market share of around 60% [1]. Especially, the so called “high performance mc-silicon” or “HPM silicon” which is commercially available since 2011 [2,3], is still competitive against the mono-crystalline Czochralski (CZ) silicon. The HPM silicon is characterized by its fine grain structure and a relatively high fraction of random grain boundaries of about 70% prohibiting the dislocation movement through the ingot volume. However, it was recently shown [4] that the benefit of the HPM material versus the classical mc material (coarse grain structure with high amount of twins) decreases with increasing ingot height. This is correlated to the strong decrease of the random (R) grain boundary fraction during growth – especially in the first centimeters – which is due to annihilation mechanisms (e.g. R + R → R, meaning two R grain boundaries combine to one single R
⁎
grain boundary) or due the formation of new low energy Σ3 twin boundaries (e.g. R → Σ3 + R, meaning a R grain boundary splits into another R grain boundary plus a new twin boundary) [4,5] as shown in Fig. 1. In order to further improve the grain structure properties of HPM silicon, two approaches concerning the random grain boundary fraction could be helpful and were evaluated within this work. The further increase of the initial random grain boundary fraction (marked with “1” in Fig. 1) might suppress the development of dislocations over a larger ingot height. Secondly, the same effect could be achieved by preferably keeping the R grain boundary fraction as high as possible along the complete ingot height (marked with “2” in Fig. 1). It was previously shown that the initial grain size and also the initial R grain boundary fraction at the bottom of classical silicon seeded HPM ingots depends on the internal microstructure of the used seeding feedstock [6]. The mean grain size within two typical feedstock variants from Siemens process (SIE) and fluidized bed reactor (FBR) lays in the
Corresponding author at: Fraunhofer Institute IISB, Schottkystr. 10, 91058 Erlangen, Germany. E-mail address: [email protected] (M. Trempa).
https://doi.org/10.1016/j.jcrysgro.2019.03.005 Received 14 November 2018; Received in revised form 4 March 2019; Accepted 5 March 2019 Available online 06 March 2019 0022-0248/ © 2019 Elsevier B.V. All rights reserved.
Journal of Crystal Growth 514 (2019) 114–123
M. Trempa, et al.
dimensions of 220 × 220 × 130 mm3 and a weight of 15 kg. In all cases, standard fused silica crucibles with a standard Si3N4 release coating and polysilicon feedstock from Wacker were used. In total three experiment series were carried out: – Series I addresses the question whether a finer silicon seeding feedstock in the 0.3–3 µm range or a more closely packed layer of feedstock particles result in an even higher R grain boundary fraction in comparison to a reference ingot seeded on standard Siemens feedstock chunks, especially at the bottom of the ingot. – Series II and III address the evolution of the R grain boundaries over ingot height. Within series II, the effect of the seeding procedure (elongation direction of the initially formed grain boundaries) and within series III the influence of the growth rate was investigated. Experiment series I consists of three G1 HPM silicon ingots grown by the classical HPM seeding approach. Two types of silicon powders with D50 values of 3.8 µm and 0.3 µm, respectively, were used as silicon feedstock. This is at least two orders of magnitude smaller than the internal grain size of the typically used Siemens (70–270 µm) or FBR (700 µm) feedstock [7]. In a third experiment, a vertical cut from a Siemens rod was used in similar way as recently reported for an in industrial G6 setup [12]. In this case, the microstructure is comparable to that of the commonly used Siemens chunks, however, the single feedstock particles within the rod are more closely packed in comparison to a porous chunk layer. After crystallization, horizontal cut wafers from different ingot heights were prepared and the mean grain size as well the grain boundary type distribution were measured by the grain detector/Laue scanner system [13]. For the experimental series II, different special seed arrangements were used in G1 experiments to investigate the evolution of R grain boundaries in dependence on their geometrical arrangement. In a first step, vertically elongated grain boundaries at the ingot bottom were induced by using mono-crystalline seed stripes oriented in 〈1 1 0〉 and 〈4 2 1〉 vertical growth direction in a similar way as it is done within the “grain boundary engineering” for the growth of Quasimono crystals [14]. The seed stripes, each of them 10 mm in width, 150 mm in length and 20 mm in thickness, are placed side-byside at the crucible bottom (compare Fig. 2). According to the seed orientations in horizontal x- and y-direction on the right ingot half, several in parallel arranged R grain boundaries should be formed, whereas on the left side Σ3 grain boundaries were intended to be generated for comparison. After crystal growth, vertical cuts were prepared to investigate the grain boundary evolution by visual inspection as well to determine their type by Laue measurements. In two further G1 experiments, low symmetry Σ29 grain boundaries, meaning quite comparable to R ones, between the seed stripes were generated. In this case, the distance between two adjacent grain boundaries was reduced by using thinner 〈1 0 0〉 seed stripes (2 and 4 mm in width) which were placed between wide 〈1 1 0〉 stripes. In a similar experiment, the thin 〈1 0 0〉 seed stripes were replaced by 〈1 1 1〉 ones in order to investigate the influence of the surface energy of the grains next to the grain boundary on its evolution over ingot height. Due to crystallographic reasons, the resulting grain boundary in this case was close to a Σ9 grain boundary. The growth behavior of the intentionally induced grain boundaries was investigated on vertical cuts by PL imaging and Laue measurements. Finally, the experiment series III consists of several HPM silicon ingots with a variation of the growth rate from 3 mm/h up to 30 mm/h by adjusting the gas flow in the full-faced cooling unit just below the crucible bottom in the range from 0 m3/h to 50 m3/h (see Table 1). In a first growth run for each parameter variation, the real growth rate was measured by the mechanical dipping procedure. During further runs, silicon ingots were grown under identical growth conditions and subsequently analyzed. As seed layer for these HPM ingots, silicon feedstock chunks (SIE) as well as 20 mm thick horizontal cuts from prior
Fig. 1. Grain boundary length fraction for Random and Σ3 grain boundaries in HPM silicon over ingot height. Additionally, two possible optimization approaches are marked with “1” and “2”.
range of 70–270 µm and 700 µm, respectively [7]. Therefore, in this work it is investigated whether a silicon seeding feedstock with a finer internal grain size of ≪70 µm is leading to an even smaller grain size and whether an increase of the initial R grain boundary fraction is achievable. Further, it is assumed that the frequently observed annihilation process of the R grain boundaries is strongly supported by the mostly diagonally elongated grain boundaries in the initial state of growth which are predetermined by the irregularly arranged Si feedstock particles (classical HPM approach) or induced by the randomly orientated nuclei at the surface of the used functional coatings (HPM 2.0). Therefore, it was investigated how the initial geometrical alignment of the R grain boundaries affects the annihilation process during the first centimeters of grain growth. This was checked by growth experiments in G1 scale where special silicon seed arrangements were used in order to predetermine the geometrical direction of the initially formed R grain boundaries. Further possible factors of influence on the grain boundary behavior and especially on the annihilation of R grain boundaries could be process parameters like the axial temperature gradient in the crystallization furnace or the growth rate of the HPM silicon ingot itself. So far, there exist some hints about that in the literature, but the proposals are contradictory. Wong et al. [8] and Lin et al. [9] investigated small mc Si ingots and EMC mc Si wafers, respectively, and observed that the higher the growth rate the faster the R grain boundary fraction decreases with ingot height due to the increase of newly formed Σ3 twin boundaries. On the contrary, Chen et al. [10] recently showed that a higher growth rate, induced by a larger axial temperature gradient, during growth of industrial G6 ingots can lead to a more columnar grain structure preventing the interaction of grain boundaries and therefore attenuating the decrease of the R grain boundary fraction over ingot height. In order to clarify this contradiction, several G1 experiments were carried out with the growth rate varied between 3 mm/h and 30 mm/h in comparison to the reference case of 10 mm/h already presented in a previous work [4].
2. Experimental setup and characterization All crystallization experiments were carried out in a G1 crystallization furnace [11] allowing the growth of G1 silicon ingots with 115
Journal of Crystal Growth 514 (2019) 114–123
M. Trempa, et al.
Fig. 2. Seed arrangement for inducing Σ3 twin boundaries and Random (R) grain boundaries on the left and right ingot half, respectively.
occur due to statistical reasons, because only a fraction of the grain boundaries can be measured for such fine grain structures (for better visibility the error bars are shown only for two data points per ingot height). As shown in Fig. 4 the resulting mean grain area for the use of the 3.8 µm Si-powder and the SIE rod is comparable to the three reference cases (seeding on SIE chunks), whereas the use of the 0.3 µm powder leads to a slightly higher value at 5 mm grown ingot height. At 80 mm ingot height the values for the new seeding approaches lay within the range of the three reference ingots. This confirms the optical impression from the grain structure images in Fig. 3. Concerning the grain boundary length fraction, no clear difference to the references is observable within the measurement accuracy. However, at 20 mm ingot height, a slight decrease of the R grain boundary fraction of 3–5% relative to the reference crystals is visible for all data points of the new seeding approaches. This means that a reduction of the inner microstructure size from ∼70–200 µm in the SIE chunks down to some microns or even to the nanometer scale has no significant effect on grain structure properties at the investigated ingot heights of ≥5 mm. Furthermore, a more closely packed layer of chunks, which is simulated by the use of the SIE rod, shows no significant effect. This is in contradiction to the initial assumption that a finer microstructure should lead to an increased number of grains and also an increased R grain boundary fraction due to the higher number of seeding sites at the initial state. To shed light on this contradiction, the first mm of growth, which cannot meaningfully be measured by the grain analysis tools due to the too small grain size, were additionally investigated by light microscopy on mechanically polished vertical cuts which were etched for 2 min in a Secco solution. In Fig. 5 the seed level regions of one reference ingot (seeded on SIE chunks) and of the ingots seeded on the 0.3 µm Si particle layer and on the SIE rod are shown. In the case of the reference the grain sizes close to the macroscopic observable seeding interface (red line) vary in a wide range. Directly on top of a Si chunk very small grains are formed, while in the regions away from the chunks the grain size is larger (marked by circles). This is mainly due to the porous chunk layer structure and the infiltrated and solidified Si melt. Further, it can be observed that the small, newly grown grains in the vicinity of the chunks are existent for a few hundred microns of grown ingot height only until the grain selection results in coarser grains. A similar effect is observable for the ingot seeded on the very fine
Table 1 Overview about growth conditions of G1 experiments within experiment series III. Experiment
Cooling gas flow [m3/h]
Measured growth rate [mm/h]
“Low” “Reference” “High” “Extremely High”
0 10 30 50
2–6 2–10 9–18 12–30
experiments (explanation see later in Section 3) were used. After crystallization, the silicon ingots were squared to 156 × 156 mm2 bricks and horizontal wafers at different ingot heights were prepared for grain structure analysis (procedure see [4,13]) and photoluminescence imaging (procedure see [4]) in order to determine the grain boundary type distribution and the area of recombination active dislocation clusters. 3. Results and discussion 3.1. Series I: Influence of the seeding material on the grain boundary distribution In Fig. 3 the grain structure images for the G1 ingots seeded on different Si materials (experiment series I) are shown for ingot heights at 5 mm and 20 mm above the seeding interface. In case of the reference, meaning the seed layer consists of small silicon chunks (Wacker polysilicon size 1, 3–15 mm particle size) from the Siemens process, two identical processed ingots were shown to obtain some statistics about the experimental deviation. Already the optical impression gives the hint that the grain structure properties regarding the grain size are not very different for all grown ingots. Fig. 4 shows the mean grain area and the R grain boundary length fraction over ingot height. The mean grain area is calculated by sum-up all grain areas divided by the total number of grains. Typically, the mean grain area for our conventionally grown G1 HPM silicon ingots starts with ∼1 mm2 at 5 mm grown ingot height and increases up to ∼3.5–4 mm2 at 80 mm. The measurement accuracy here is in the range of 0.05 mm2 in the lower ingot parts and 0.1 mm2 in the upper regions. The random grain boundary fraction typically begins with ∼75% at 20 mm and drops down to 62–67% at 80 mm. An error of up to 5% can 116
Journal of Crystal Growth 514 (2019) 114–123
M. Trempa, et al.
Fig. 3. Grain structure of HPM-Si ingots seeded on Siemens chunks (two references), a cut from a Siemens rod and two different Si-powders (3.8 µm and 0.3 µm particle size). The images were taken at 5 mm and 20 mm above the particular seeding level. For better visibility a magnification of a 20 × 20 mm2 areas is shown for 5 mm ingot height.
(marked with lower circle). Therefore the mean grain size directly above the seeding interface is likely to be smaller than for the other ingots. However, also here the grain size strongly increases within the first several hundred microns caused by grain selection, resulting in a comparable grain size distribution at larger ingot heights (marked with upper circle). In summary, even though alternative seeds might yield a finer initial grain structure, grain selection results in a quick increase of the grain size within the first few hundred microns of growth. This results in comparable grain structure properties at relevant ingot heights ≥5 mm. In conclusion it is not possible to significantly reduce the mean grain
0.3 µm Si-powder. While the grains in the vicinity of the Si-particles are very small, larger grains are found in the cavities of the powder layer (marked with circles). Also here grain selection leads to a quick coarsening of the initially fine grain structure. Additionally, the fine powder particles obviously agglomerate to larger particles which reduce the assumed advantage of the powder providing smaller seeding crystals. Finally, the ingot seeded on the SIE rod shows the effect of grain coarsening due to grain selection even more clearly. In this case, due to the lack of the porosity of the seeding layer and therefore no existing infiltration of the Si melt into the seed layer, the fine grains were continuously formed along the complete area of the seeding interface
117
Journal of Crystal Growth 514 (2019) 114–123
M. Trempa, et al.
4,5
R grain boundary length fraction [%]
85
Mean grain area [mm²]
4,0 3,5 3,0 2,5 2,0
SIE chunks SIE chunks (rep.1) SIE chunks (rep.2) SIE rod 3.8μm Si-powder 0.3μm Si-powder
1,5 1,0 0,5
80
75
70
65
60
55
0
20
40
60
80
10
Grown ingoth height [mm]
20
30
40
50
60
70
80
Grown ingot height [mm]
Fig. 4. Mean grain area (left) and grain boundary length fraction (right) at different grown ingot heights for HPM Si ingots seeded on Siemens chunks (3×), a cut from a Siemens rod and two different Si-powders.
Fig. 5. Light microscope images of polished and Secco-etched vertical cuts in the seed-level region of different HPM Si ingots: Reference grown on SIE chunks (top), SIE rod (middle) and 0.3 µm Si-powder (bottom). 118
Journal of Crystal Growth 514 (2019) 114–123
M. Trempa, et al.
Fig. 6. Vertical cut out of the center of a G1 ingot with artificially induced grain boundaries: (a) grain boundary types, (b) crystal orientation parallel to growth direction and (c) growth interface shape measured by lateral photovoltage scanning (LPS). The green arrows indicate grain boundary growth perpendicular to the growth interface, the red arrows a non-perpendicular growth.
energy minimization causing a roof-like structure above the seed stripes [15]. The fact that one roof side is labeled by the Laue-Scanner software as a R grain boundary is according to the used Brandon-criterion [16] justifying the classification of the grain boundaries, although it has to be a Σ3 or Σ9 type due to the crystallographic relationship of the grains in this region. These non-perfect twin boundaries slightly exhibit the tolerable angle deviations larger than 8.7° (in the case of Σ3) and 5° (in the case of Σ9) which results in the labeling as a R grain boundary. Nevertheless, all new formed twin boundaries grow straightly upwards according to the direction of their twin plane and not according to the shape of the solid liquid interface, as can be seen in Fig. 6c and is expected for low index Σ grain boundaries. On the right half of the ingot, seven parallel R grain boundaries were formed by the seeding method. It is obvious that at the beginning of solidification they are growing more or less parallel to each other, however a slight tilt between the grain boundary lines is observable with increasing ingot height. This can be well correlated to the growth interface shape in this ingot region meaning the R grain boundaries propagate perpendicularly to the slightly curved solid-liquid interface as expected (green arrows in Fig. 6c). At about half of the ingot height, two pairs of neighboring R grain boundaries abruptly align each other and coincide with the result that the number of grain boundaries is reduced from four random grain boundaries to two small angle grain boundaries which are not visible in this image. It has to be noticed that in most cases in real HPM growth experiment the annihilation reaction would be R + R → R due to the generally non-equal grain orientation of the involved, next-neighboring grains. This annihilation phenomenon cannot be explained only by a direct correlation to the almost flat shape of the solid-liquid interface, as can be seen by the red arrows in Fig. 6c. In this case the growth direction of the grain boundaries deviates in some degrees from the perpendicular direction to the growth interface, indicating that the
size or increase the random grain boundary fraction at 5 mm ingot height and above in comparison to the reference HPM ingots by the seeding variants presented above. 3.2. Series II: Influence of seeding arrangement on the evolution of R grain boundaries during growth The aim of the experiment series II was to investigate if it is possible to keep alive R grain boundaries in HPM material over ingot height if they were arranged parallel to each other and perpendicular to the crucible bottom, respectively. For that purpose alternating 〈1 1 0〉 and 〈4 2 1〉 oriented seeds were used (arrangement shown in Fig. 2). Because it is well known that the growth interface shape can strongly influence the growth direction of non-twin grain boundaries [17], the growth interface should be as flat as possible. The illustration of the measured growth interface in Fig. 6 c shows an almost flat interface with a slight W-shaped character in the upper part of the ingot. Unfortunately, the exact shape of the growth interface could not be detected in the lower ingot part due to the electrical properties of the silicon seeds. Because the interface becomes more flat towards the top, it can be assumed that the W-shape in the lower part is little bit more pronounced. A vertical cut from the center of this special seeded ingot was investigated by the Laue method. The resulting grain boundary and grain orientation map are shown in Fig. 6a/b. The 10 mm wide seed stripes used for the grain boundary engineering are visible below the seeding level at the bottom of these maps. The ingot can be divided in two parts: In the left half, Σ3 twin boundaries were intentionally induced whereas in the right half, R grain boundaries were generated. Just above the seeding interface, the initially formed Σ3 twin boundaries split into new straightly aligned twin boundaries (examples are marked with white circles), due to 119
Journal of Crystal Growth 514 (2019) 114–123
M. Trempa, et al.
Fig. 7. Top view of seed arrangements and PL-images of vertical cuts of the two G1 ingots grown with thin 〈1 0 0〉 (left) and 〈1 1 1〉 (right) seed stripes. The PLimages are made up of different PL-measurements in order to make the grain boundaries visible over the complete ingot height.
overgrowth of the stripe between the two grain boundaries is triggered by the surface energy of the adjacent grains. In order to validate this hypothesis of competing crystal regions, a further G1 experiment was carried out for which the 〈1 0 0〉 stripes were replaced by 〈1 1 1〉 oriented stripes with 2 mm thickness (see Fig. 7 on the right). The {1 1 1} surface of silicon is known to be the surface with the lowest surface energy. Indeed, the 〈1 1 1〉 oriented stripe region becomes wider with increasing ingot height meaning that the Σ9 grain boundaries are further separated. Interestingly, at ca. 50 mm growth height, the right grain boundary exhibits a kink towards the left grain boundary (marked by the lower red arrow). Laue measurements reveal that this is caused by a twinning process which is well known from the Quasimono technology using monocrystalline seed plates [15]. In this case, the orientation relationship between the 〈1 1 0〉 seeds and the 〈1 1 1〉 seed stripe promotes the twin formation at the edge of the grown 〈1 1 1〉 grain which results in formation of new (twin) boundaries. The initially formed Σ9 grain boundaries are replaced by two Σ3 grain boundaries growing further in upward direction. At approx.70 mm in ingot height the twinning process reoccurs. Again two new Σ3 grain boundaries are formed growing until the top of the ingot. In conclusion, the evolution of the random grain boundaries and all other non-twin boundaries is mainly influenced by the shape of the
kinetics of the surface energy of the adjacent grains plays a more crucial role [18]. Further, it is well known that grains overgrow others with higher surface energy, provided that no strong other influences like highly curved growth interfaces or high growth rates are existent [19]. In the present case, the 〈1 1 0〉 orientation would then have a lower surface energy than the 〈4 2 1〉 orientation (compare grain orientation map in Fig. 6b). Additionally, it can be remarked that the meeting of the grain boundaries in the shown 2D cut is a consequence of other (similar) grain boundary interactions which takes place at lower ingot heights, but at other radial position. A similar observation was made for the other G1 experiment of series II with thin 〈1 0 0〉 oriented seed stripes positioned between 〈1 1 0〉 oriented seed plates (see Fig. 7 on the left). On PL-images of a vertical cut through the ingot center it is visible that both of the left pair of generated Σ29 grain boundaries (type was confirmed by Laue measurements) are grown roughly parallel to each other in the lower part of the ingot up to a certain point at which they were annihilated by overgrowth of the center grain and formation of a small angle grain boundary (SA-GB). On the right, the setup was identical, but the distance between the two grain boundaries was twice as large, i.e. 4 mm instead of 2 mm. This resulted in longer Σ29 grain boundaries which are however also annihilated by overgrowth. The corresponding positions are marked with yellow circles. Also here it is probable that the 120
Journal of Crystal Growth 514 (2019) 114–123
M. Trempa, et al.
h), the growth rate is almost identical at below 5 mm/h at the early beginning of growth which is caused by the thermal inertia of the furnace setup. For the high gas flow rates, differences are already present at the beginning where the growth rate is increased to 9 mm/h and 13 mm/h for 30 m3/h and 50 m3/h gas flow, respectively. The experiment with the cooling gas flow of 10 m3/h equals to the process applied for the growth of a 710 mm high ingot which was simulated by the subsequent growth of eight G1 ingots published in our previous work [4]. For the first ingot, the HPM grain structure was induced by a 20 mm seeding layer containing silicon chunks. For the following seven ingots, a 20 mm thick seed plate cut from the upper part of the previously grown ingot was used. In Fig. 9 the grain boundary length fraction over the total ingot height of 710 mm is shown for the reference experiment (cooling 10 m3/h, growth rate ∼2–10 mm/h) by the squared unfilled symbols [4]. It can be observed that the R grain boundary fraction strongly decreases from 73% to 50% along the first 150 mm grown ingot height and remains constant up to 710 mm, which was interpreted as stable state with annihilation and formation processes being in equilibrium [4]. It has to be noticed that this found equilibrium state is probably linked to the appearing growth conditions in our G1 furnace. In the case of industrial setups the exact values of the grain boundary plateau and the ingot height at which the equilibrium is reached can be different. In the present work we investigated whether a different growth rate would change the grain boundary type distribution over ingot height, especially along the first 150 mm ingot height and within the plateau at ingot heights over 400 mm. For testing the effect of the decreased growth rate of 2–6 mm/h in the lower ingot region, one G1 ingot was grown using silicon chunks as seeds, i.e. starting at an initial ingot height of 0 mm. The effect in the plateau region at this growth rate was investigated by first growing an ingot with a 30 mm thick slice cut from the reference ingot at 370–400 mm height as seed plate, yielding data points at 400 mm and 460 mm. The top of this ingot was then again used as seed for growth of another ingot, allowing to obtain data for a virtual ingot height of 560 mm. The results in Fig. 9 (filled triangle symbols) clearly show that the lower growth rate has no measurable influence on the R grain boundary fraction, both in the lower ingot region as well as in the stable state region. The effect of the higher growth rate of 9–18 mm/h was investigated in a similar way. In this case, three subsequent ingots were grown with the first starting from the silicon feedstock, i.e. from an ingot height of 0 mm. Additionally, one ingot was grown within the stable state region by using a seed plate cut from the reference ingot at 460–490 mm grown ingot height. The results are also plotted in Fig. 9 (filled circle symbols). Within the measurement accuracy again no clear difference to the reference is observable. However, in the first 240 mm ingot height, for all data points a slight decrease of the R grain boundary fraction and a slight increase of twins of 1–3% relative to the trend line of the reference crystal is visible. This trend would agree to the observations of Wong et al. [8] and Lin et al. [9] who proposed an enhanced twin formation with increasing growth rate due to the higher probability of undercooling. Finally, one single ingot was grown with the extremely high growth rate of 12–30 mm/h starting from the silicon feedstock seeding layer (ingot height = 0 mm). Due to the highly concave curvature of the growth interface in the upper parts of the ingot, only two wafers at 35 and 45 mm ingot height could be evaluated. From these results (filled rhombic symbols), again no clear influence of the growth rate on the R grain boundary fraction was observed. Comparably to the other experiment with increased growth rate, the grain boundary type distribution could be slightly shifted to lower R and higher twin values, but again with deviations only within the margin of error. In summary, the realized variation in the maximum growth rate by a factor of 0.5 up to 3 has no significant influence on the grain boundary type distribution. At most, only a slight, but insignificant decrease of the R grain boundary fraction was observed, which we
solid-liquid interface. Additionally, the surface energy of the adjacent grains slightly influences the growth angle of the grain boundaries due to kinetical reasons. The first point is clearly visible also by the “grain boundary arrangement” within the SIE rod experiment in series I. As shown in Fig. 5, the grain boundaries within the rod are arranged in directions from diagonal to parallel with respect to the seeding interface. Despite of this structure, the grain boundaries of the newly formed grains develop more or less perpendicularly to the even flat seeding interface which means their evolution follows the growth interface shape. Based on the above described the following rules to reduce/avoid the annihilation processes of initially formed random grain boundaries can be laid down:
• The initial random grain boundaries should be arranged parallel to each other (perpendicular to the seeding interface). • The distance between neighboring random grain boundaries • • •
shouldn’t be too small in order to shift the meeting of them to larger ingot heights. The growth interface shape should be as flat as possible. The orientations of the adjacent grains should be similar regarding their surface energy. The orientation relationship between adjacent grains should not promote twinning processes.
All these points, with exception of the flat growth interface, are very difficult to achieve within an industrial growth setup. In particular, some of the mentioned points are in contradiction to the use of randomly oriented seeds as it is applied for the HPM process. Even though artificially arranged seeding layers yielding a high and stable random grain boundary fraction might exist, the required seed arrangement would be even more complex than for the Quasi-Mono approach. Therefore, this approach is probably no realistic option from an economic point of view. 3.3. Series III: Influence of growth rate on the growth behavior of R grain boundaries To investigate the influence of the growth rate, four G1 ingots were crystallized with argon cooling gas flow rates of 0 m3/h, 10 m3/h, 30 m3/h and 50 m3/h. The resulting growth rates for these conditions are depicted in Fig. 8. As intended, an increasing gas flow results in a higher growth rate. In the case of low gas flow rates (0 m3/h and 10 m3/
Growth rate [mm/h]
40
50 m³/h 30 m³/h 10 m³/h 0 m³/h
30
20
10
0 0
20
40
60
80
100
120
Ingot height [mm] Fig. 8. Measured growth rate over the grown ingot height for varying cooling gas flow conditions. 121
Journal of Crystal Growth 514 (2019) 114–123
grain boundary length fraction [%]
M. Trempa, et al.
80
R Σ3
70 60
2-6 mm/h
50
2-10 mm/h (Ref.)
40
9-18 mm/h
30
12-30 mm/h
20 10 0
0
50
100
200 300 400 500 600 700
grown ingot height [mm] Fig. 9. Grain boundary length fraction over the grown ingot height for different growth rates.
growth rate of 2–6 mm/h at the ingot heights 460 mm and 560 mm. This difference is also obvious to the naked eye from the PL images on the right hand side in Fig. 10. Since the initial dislocated area in these ingots should be about 12% (the seed plate was cut at 370–400 mm from the reference ingot), the dislocated area is decreased down to 8%. This indicates that some of the existing dislocation clusters were overgrown by dislocation free grains. Maybe also less thermal stress is generated by the lower growth rate, which reduces the dislocation multiplication.
assess as deterioration instead of an improvement relative to the reference process. Even though this might indicate that even lower growth rates than the 2–6 mm/h tested here could yield a (slightly) increased R grain boundary fraction, at least in the lower region of the ingot, such an approach would not be feasible in industry due to economic considerations. Thus, the authors suggest that the process parameter growth rate is no tool to further improve the random grain boundary fraction in industrially produced HPM silicon ingots. Even though the grain structure has been proven to be a very good indicator for the resulting solar cell performance, it is only an indirect method relying on the correlation between grain structure and emergence of dislocations. Therefore, beside the grain structure analysis, the wafers were also investigated by means of PL imaging to determine the recombination active dislocation clusters. In Fig. 10, the recombination active wafer area, occupied by sum-up the area covered by dislocation clusters (grey-value analysis) divided by the total wafer area, is plotted along the grown ingot height. For the highest growth rate of 12–30 mm/h no statement can be made due to the lack of samples in ingot heights > 45 mm. Further, an increased growth rate of 9–18 mm/h yields a similar trend of the recombination active area over the ingot height as the reference. In contrast, a reduced dislocated area relative to the reference is measured for the lower
4. Conclusion Several hypothetical approaches to improve the grain structure properties of HPM silicon ingots were evaluated by experiments in G1 scale. It was shown that the use of extremely fine Si particles or densely packed chunk layers can lead to a finer initial grain structure and probably also to a higher length fraction of the suitable random grain boundaries, but only in the first few 100 µm of growth. The natural grain selection results in a quick grain coarsening and therefore no positive effect remains at relevant block heights. From grain boundary engineering experiments we found some directives, like a parallel arrangement of the initial random grain
Fig. 10. Recombination active area over grown ingot height for different process conditions. On the right hand side the PL-images of three ingots at ∼550 mm ingot height are shown. 122
Journal of Crystal Growth 514 (2019) 114–123
M. Trempa, et al.
boundaries or keeping flat the growth interface, which may help to reduce the annihilation of the random grain boundaries. However, an industrial implementation seems to be rather complicated, especially concerning the even complex seed arrangement which has to fulfill a lot of requirements. Finally, also the growth rate, the influence of which was tested over a wide range, has no influence on the annihilation processes of the random grain boundaries. Nevertheless, the area fraction of electrically active dislocations was found to be reduced for applying a reduced growth rate, which again is unfortunately undesirable in industry processes for economic reasons. In summary, it has to be recognized that from the grain structure point of view, the properties of the current HPM process are most probably the best ones which can be achieved, at least under industrial conditions, and therefore further improvements of HPM material quality has to be approached via other paths, e.g. the reduction of thermal stress or of impurities like metals or oxygen.
References [1] ITRPV, 2018. International Technology Roadmap for Photovoltaic (ITRPV), 2017. ninth ed. [2] C.W. Lan, W.C. Lan, T.F. Lee, A. Yu, Y.M. Yang, W.C. Hsu, B. Hsu, A. Yang, J. Crys. Growth 360 (2012) 68–75. [3] Y.M. Yang, A. Yu, B. Hsu, W.C. Hsu, A. Yang, C.W. Lan, Prog. Photovoltaics. Res. Appl. 23 (3) (2015) 340–351. [4] M. Trempa, I. Kupka, C. Kranert, T. Lehmann, C. Reimann, J. Friedrich, J. Crys. Growth 459 (2017) 67–75. [5] R.R. Prakash, K. Jiptner, J. Chen, Y. Miyamura, H. Harada, T. Sekiguchi, Appl. Phys. Express 8 (3) (2015) 35502. [6] C. Reimann, M. Trempa, T. Lehmann, K. Rosshirt, J. Stenzenberger, J. Friedrich, K. Hesse, E. Dornberger, J. Crys. Growth 434 (2016) 88–95. [7] K.E. Ekstrøm, G. Stokkan, A. Autruffe, R. Søndenå, H. Dalaker, L. Arnberg, M. Di Sabatino, J. Crys. Growth 441 (2016) 95–100. [8] Y.T. Wong, C. Hsu, C.W. Lan, J. Crys. Growth 387 (2014) 10–15. [9] H.K. Lin, M.C. Wu, C.C. Chen, C.W. Lan, J. Crys. Growth 439 (2016) 40–46. [10] W. Chen, Q. Wang, D. Yang, L.D. Li, X.G. Yu, L. Wang, H. Jin, J. Crys. Growth 467 (2017) 65–70. [11] M. Beier, M. Trempa, J. Seebeck, C. Reimann, P. Wellmann, B. Gründig-Wendrock, J. Friedrich, K. Dadzis, L. Sylla, T. Richter, in: Proc. of the 29th European Photovoltaic Solar Energy Conference, Amsterdam, pp. 717–721. [12] C. Huang, H. Zhang, S. Yuan, Y. Wu, X. Zhang, D. You, L. Wang, X. Yu, Y. Wan, D. Yang, Sol. Energ. Mat. Sol. C. 179 (2018) 312–318. [13] T. Lehmann, M. Trempa, E. Meissner, M. Zschorsch, C. Reimann, J. Friedrich, Acta Mater. 69 (2014) 1–8. [14] M. Trempa, C. Reimann, J. Friedrich, G. Müller, A. Krause, L. Sylla, T. Richter, Cryst. Res. Technol. 50 (2015) 124–132. [15] M. Trempa, C. Reimann, J. Friedrich, G. Müller, D. Oriwol, J. Crys. Growth 351 (2012) 131–140. [16] D.G. Brandon, Acta Metall. 14 (11) (1966) 1479–1484. [17] Klaus-Thomas Wilke, Joachim Bohm (Hg.), Kristallzüchtung. 1. Aufl. [18] T. Duffar, A. Nadri, Crystal Growth/Croissance Cristalline 14 (2) (2013) 185–191. [19] K. Fujiwara, Y. Obinata, T. Ujihara, N. Usami, G. Sazaki, K. Nakakjima, J. Crys. Growth 266 (2004) 441–448.
Acknowledgement These research and development activities are partly carried out within the ENOWA-II project (0325805E) funded by the German Federal Ministry for Economic Affairs and Energy. Appendix A. Supplementary material Supplementary data to this article can be found online at https:// doi.org/10.1016/j.jcrysgro.2019.03.005.
123
Contents lists available at ScienceDirect
Journal of Crystal Growth journal homepage: www.elsevier.com/locate/jcrysgro
Evaluation of improvement strategies of grain structure properties in high performance multi-crystalline silicon ingots
T
⁎
M. Trempaa, , C. Kranerta,b, I. Kupkaa,b, C. Reimanna,b, J. Friedricha,b a b
Fraunhofer IISB, Schottkystr. 10, 91058 Erlangen, Germany Fraunhofer THM, Am St.-Niclas-Schacht 13, 09599 Freiberg, Germany
A R T I C LE I N FO
A B S T R A C T
Communicated by A.G. Ostrogorsky
High performance multi-crystalline silicon (HPM-Si) for the use in photovoltaics is characterized by a very fine grain structure and a high content of random grain boundaries, finally resulting in a low dislocation density and consequently in a high material quality. Typically, the grain size increases and the fraction of random grain boundaries decreases over ingot height due to annihilation mechanisms, especially in the first 150 mm. One approach for further material improvement is to further increase the initial random grain boundary fraction and to maintain it as high as possible over the complete ingot height. In this work, several theoretical approaches to achieve these points were evaluated by experiments in G1 scale. Firstly, the influence of the silicon seeding material on the initial grain structure was investigated regarding the effect of extremely fine Si particles in the µm to nm range and the bulk density of the particle layer. Secondly, the effect of the initial geometrical grain boundary arrangement in the seed layer was evaluated. For that purpose, special seed alignments similar to the Quasimono approach were tested. Finally, the process parameter growth rate was varied in a wide range to investigate its influence on the evolution of the grain boundaries during growth. The results show that the optimum for the initial random grain boundary fraction is already reached by the existing methods/commonly used seed materials. Concerning the decrease of the random grain boundary fraction over ingot height, some technical aspects were identified which are able to keep the amount of random grain boundaries at a high level. However, the practical realization within an industrial setup might be difficult.
Keywords: A1. Directional solidification A1. High performance multi-crystalline silicon A1. Grain structure A1. Grain boundaries A1. Recombination active defects
1. Introduction Multi-crystalline (mc) silicon produced by directional solidification technique is still the most used material for the production of silicon solar cells with a market share of around 60% [1]. Especially, the so called “high performance mc-silicon” or “HPM silicon” which is commercially available since 2011 [2,3], is still competitive against the mono-crystalline Czochralski (CZ) silicon. The HPM silicon is characterized by its fine grain structure and a relatively high fraction of random grain boundaries of about 70% prohibiting the dislocation movement through the ingot volume. However, it was recently shown [4] that the benefit of the HPM material versus the classical mc material (coarse grain structure with high amount of twins) decreases with increasing ingot height. This is correlated to the strong decrease of the random (R) grain boundary fraction during growth – especially in the first centimeters – which is due to annihilation mechanisms (e.g. R + R → R, meaning two R grain boundaries combine to one single R
⁎
grain boundary) or due the formation of new low energy Σ3 twin boundaries (e.g. R → Σ3 + R, meaning a R grain boundary splits into another R grain boundary plus a new twin boundary) [4,5] as shown in Fig. 1. In order to further improve the grain structure properties of HPM silicon, two approaches concerning the random grain boundary fraction could be helpful and were evaluated within this work. The further increase of the initial random grain boundary fraction (marked with “1” in Fig. 1) might suppress the development of dislocations over a larger ingot height. Secondly, the same effect could be achieved by preferably keeping the R grain boundary fraction as high as possible along the complete ingot height (marked with “2” in Fig. 1). It was previously shown that the initial grain size and also the initial R grain boundary fraction at the bottom of classical silicon seeded HPM ingots depends on the internal microstructure of the used seeding feedstock [6]. The mean grain size within two typical feedstock variants from Siemens process (SIE) and fluidized bed reactor (FBR) lays in the
Corresponding author at: Fraunhofer Institute IISB, Schottkystr. 10, 91058 Erlangen, Germany. E-mail address: [email protected] (M. Trempa).
https://doi.org/10.1016/j.jcrysgro.2019.03.005 Received 14 November 2018; Received in revised form 4 March 2019; Accepted 5 March 2019 Available online 06 March 2019 0022-0248/ © 2019 Elsevier B.V. All rights reserved.
Journal of Crystal Growth 514 (2019) 114–123
M. Trempa, et al.
dimensions of 220 × 220 × 130 mm3 and a weight of 15 kg. In all cases, standard fused silica crucibles with a standard Si3N4 release coating and polysilicon feedstock from Wacker were used. In total three experiment series were carried out: – Series I addresses the question whether a finer silicon seeding feedstock in the 0.3–3 µm range or a more closely packed layer of feedstock particles result in an even higher R grain boundary fraction in comparison to a reference ingot seeded on standard Siemens feedstock chunks, especially at the bottom of the ingot. – Series II and III address the evolution of the R grain boundaries over ingot height. Within series II, the effect of the seeding procedure (elongation direction of the initially formed grain boundaries) and within series III the influence of the growth rate was investigated. Experiment series I consists of three G1 HPM silicon ingots grown by the classical HPM seeding approach. Two types of silicon powders with D50 values of 3.8 µm and 0.3 µm, respectively, were used as silicon feedstock. This is at least two orders of magnitude smaller than the internal grain size of the typically used Siemens (70–270 µm) or FBR (700 µm) feedstock [7]. In a third experiment, a vertical cut from a Siemens rod was used in similar way as recently reported for an in industrial G6 setup [12]. In this case, the microstructure is comparable to that of the commonly used Siemens chunks, however, the single feedstock particles within the rod are more closely packed in comparison to a porous chunk layer. After crystallization, horizontal cut wafers from different ingot heights were prepared and the mean grain size as well the grain boundary type distribution were measured by the grain detector/Laue scanner system [13]. For the experimental series II, different special seed arrangements were used in G1 experiments to investigate the evolution of R grain boundaries in dependence on their geometrical arrangement. In a first step, vertically elongated grain boundaries at the ingot bottom were induced by using mono-crystalline seed stripes oriented in 〈1 1 0〉 and 〈4 2 1〉 vertical growth direction in a similar way as it is done within the “grain boundary engineering” for the growth of Quasimono crystals [14]. The seed stripes, each of them 10 mm in width, 150 mm in length and 20 mm in thickness, are placed side-byside at the crucible bottom (compare Fig. 2). According to the seed orientations in horizontal x- and y-direction on the right ingot half, several in parallel arranged R grain boundaries should be formed, whereas on the left side Σ3 grain boundaries were intended to be generated for comparison. After crystal growth, vertical cuts were prepared to investigate the grain boundary evolution by visual inspection as well to determine their type by Laue measurements. In two further G1 experiments, low symmetry Σ29 grain boundaries, meaning quite comparable to R ones, between the seed stripes were generated. In this case, the distance between two adjacent grain boundaries was reduced by using thinner 〈1 0 0〉 seed stripes (2 and 4 mm in width) which were placed between wide 〈1 1 0〉 stripes. In a similar experiment, the thin 〈1 0 0〉 seed stripes were replaced by 〈1 1 1〉 ones in order to investigate the influence of the surface energy of the grains next to the grain boundary on its evolution over ingot height. Due to crystallographic reasons, the resulting grain boundary in this case was close to a Σ9 grain boundary. The growth behavior of the intentionally induced grain boundaries was investigated on vertical cuts by PL imaging and Laue measurements. Finally, the experiment series III consists of several HPM silicon ingots with a variation of the growth rate from 3 mm/h up to 30 mm/h by adjusting the gas flow in the full-faced cooling unit just below the crucible bottom in the range from 0 m3/h to 50 m3/h (see Table 1). In a first growth run for each parameter variation, the real growth rate was measured by the mechanical dipping procedure. During further runs, silicon ingots were grown under identical growth conditions and subsequently analyzed. As seed layer for these HPM ingots, silicon feedstock chunks (SIE) as well as 20 mm thick horizontal cuts from prior
Fig. 1. Grain boundary length fraction for Random and Σ3 grain boundaries in HPM silicon over ingot height. Additionally, two possible optimization approaches are marked with “1” and “2”.
range of 70–270 µm and 700 µm, respectively [7]. Therefore, in this work it is investigated whether a silicon seeding feedstock with a finer internal grain size of ≪70 µm is leading to an even smaller grain size and whether an increase of the initial R grain boundary fraction is achievable. Further, it is assumed that the frequently observed annihilation process of the R grain boundaries is strongly supported by the mostly diagonally elongated grain boundaries in the initial state of growth which are predetermined by the irregularly arranged Si feedstock particles (classical HPM approach) or induced by the randomly orientated nuclei at the surface of the used functional coatings (HPM 2.0). Therefore, it was investigated how the initial geometrical alignment of the R grain boundaries affects the annihilation process during the first centimeters of grain growth. This was checked by growth experiments in G1 scale where special silicon seed arrangements were used in order to predetermine the geometrical direction of the initially formed R grain boundaries. Further possible factors of influence on the grain boundary behavior and especially on the annihilation of R grain boundaries could be process parameters like the axial temperature gradient in the crystallization furnace or the growth rate of the HPM silicon ingot itself. So far, there exist some hints about that in the literature, but the proposals are contradictory. Wong et al. [8] and Lin et al. [9] investigated small mc Si ingots and EMC mc Si wafers, respectively, and observed that the higher the growth rate the faster the R grain boundary fraction decreases with ingot height due to the increase of newly formed Σ3 twin boundaries. On the contrary, Chen et al. [10] recently showed that a higher growth rate, induced by a larger axial temperature gradient, during growth of industrial G6 ingots can lead to a more columnar grain structure preventing the interaction of grain boundaries and therefore attenuating the decrease of the R grain boundary fraction over ingot height. In order to clarify this contradiction, several G1 experiments were carried out with the growth rate varied between 3 mm/h and 30 mm/h in comparison to the reference case of 10 mm/h already presented in a previous work [4].
2. Experimental setup and characterization All crystallization experiments were carried out in a G1 crystallization furnace [11] allowing the growth of G1 silicon ingots with 115
Journal of Crystal Growth 514 (2019) 114–123
M. Trempa, et al.
Fig. 2. Seed arrangement for inducing Σ3 twin boundaries and Random (R) grain boundaries on the left and right ingot half, respectively.
occur due to statistical reasons, because only a fraction of the grain boundaries can be measured for such fine grain structures (for better visibility the error bars are shown only for two data points per ingot height). As shown in Fig. 4 the resulting mean grain area for the use of the 3.8 µm Si-powder and the SIE rod is comparable to the three reference cases (seeding on SIE chunks), whereas the use of the 0.3 µm powder leads to a slightly higher value at 5 mm grown ingot height. At 80 mm ingot height the values for the new seeding approaches lay within the range of the three reference ingots. This confirms the optical impression from the grain structure images in Fig. 3. Concerning the grain boundary length fraction, no clear difference to the references is observable within the measurement accuracy. However, at 20 mm ingot height, a slight decrease of the R grain boundary fraction of 3–5% relative to the reference crystals is visible for all data points of the new seeding approaches. This means that a reduction of the inner microstructure size from ∼70–200 µm in the SIE chunks down to some microns or even to the nanometer scale has no significant effect on grain structure properties at the investigated ingot heights of ≥5 mm. Furthermore, a more closely packed layer of chunks, which is simulated by the use of the SIE rod, shows no significant effect. This is in contradiction to the initial assumption that a finer microstructure should lead to an increased number of grains and also an increased R grain boundary fraction due to the higher number of seeding sites at the initial state. To shed light on this contradiction, the first mm of growth, which cannot meaningfully be measured by the grain analysis tools due to the too small grain size, were additionally investigated by light microscopy on mechanically polished vertical cuts which were etched for 2 min in a Secco solution. In Fig. 5 the seed level regions of one reference ingot (seeded on SIE chunks) and of the ingots seeded on the 0.3 µm Si particle layer and on the SIE rod are shown. In the case of the reference the grain sizes close to the macroscopic observable seeding interface (red line) vary in a wide range. Directly on top of a Si chunk very small grains are formed, while in the regions away from the chunks the grain size is larger (marked by circles). This is mainly due to the porous chunk layer structure and the infiltrated and solidified Si melt. Further, it can be observed that the small, newly grown grains in the vicinity of the chunks are existent for a few hundred microns of grown ingot height only until the grain selection results in coarser grains. A similar effect is observable for the ingot seeded on the very fine
Table 1 Overview about growth conditions of G1 experiments within experiment series III. Experiment
Cooling gas flow [m3/h]
Measured growth rate [mm/h]
“Low” “Reference” “High” “Extremely High”
0 10 30 50
2–6 2–10 9–18 12–30
experiments (explanation see later in Section 3) were used. After crystallization, the silicon ingots were squared to 156 × 156 mm2 bricks and horizontal wafers at different ingot heights were prepared for grain structure analysis (procedure see [4,13]) and photoluminescence imaging (procedure see [4]) in order to determine the grain boundary type distribution and the area of recombination active dislocation clusters. 3. Results and discussion 3.1. Series I: Influence of the seeding material on the grain boundary distribution In Fig. 3 the grain structure images for the G1 ingots seeded on different Si materials (experiment series I) are shown for ingot heights at 5 mm and 20 mm above the seeding interface. In case of the reference, meaning the seed layer consists of small silicon chunks (Wacker polysilicon size 1, 3–15 mm particle size) from the Siemens process, two identical processed ingots were shown to obtain some statistics about the experimental deviation. Already the optical impression gives the hint that the grain structure properties regarding the grain size are not very different for all grown ingots. Fig. 4 shows the mean grain area and the R grain boundary length fraction over ingot height. The mean grain area is calculated by sum-up all grain areas divided by the total number of grains. Typically, the mean grain area for our conventionally grown G1 HPM silicon ingots starts with ∼1 mm2 at 5 mm grown ingot height and increases up to ∼3.5–4 mm2 at 80 mm. The measurement accuracy here is in the range of 0.05 mm2 in the lower ingot parts and 0.1 mm2 in the upper regions. The random grain boundary fraction typically begins with ∼75% at 20 mm and drops down to 62–67% at 80 mm. An error of up to 5% can 116
Journal of Crystal Growth 514 (2019) 114–123
M. Trempa, et al.
Fig. 3. Grain structure of HPM-Si ingots seeded on Siemens chunks (two references), a cut from a Siemens rod and two different Si-powders (3.8 µm and 0.3 µm particle size). The images were taken at 5 mm and 20 mm above the particular seeding level. For better visibility a magnification of a 20 × 20 mm2 areas is shown for 5 mm ingot height.
(marked with lower circle). Therefore the mean grain size directly above the seeding interface is likely to be smaller than for the other ingots. However, also here the grain size strongly increases within the first several hundred microns caused by grain selection, resulting in a comparable grain size distribution at larger ingot heights (marked with upper circle). In summary, even though alternative seeds might yield a finer initial grain structure, grain selection results in a quick increase of the grain size within the first few hundred microns of growth. This results in comparable grain structure properties at relevant ingot heights ≥5 mm. In conclusion it is not possible to significantly reduce the mean grain
0.3 µm Si-powder. While the grains in the vicinity of the Si-particles are very small, larger grains are found in the cavities of the powder layer (marked with circles). Also here grain selection leads to a quick coarsening of the initially fine grain structure. Additionally, the fine powder particles obviously agglomerate to larger particles which reduce the assumed advantage of the powder providing smaller seeding crystals. Finally, the ingot seeded on the SIE rod shows the effect of grain coarsening due to grain selection even more clearly. In this case, due to the lack of the porosity of the seeding layer and therefore no existing infiltration of the Si melt into the seed layer, the fine grains were continuously formed along the complete area of the seeding interface
117
Journal of Crystal Growth 514 (2019) 114–123
M. Trempa, et al.
4,5
R grain boundary length fraction [%]
85
Mean grain area [mm²]
4,0 3,5 3,0 2,5 2,0
SIE chunks SIE chunks (rep.1) SIE chunks (rep.2) SIE rod 3.8μm Si-powder 0.3μm Si-powder
1,5 1,0 0,5
80
75
70
65
60
55
0
20
40
60
80
10
Grown ingoth height [mm]
20
30
40
50
60
70
80
Grown ingot height [mm]
Fig. 4. Mean grain area (left) and grain boundary length fraction (right) at different grown ingot heights for HPM Si ingots seeded on Siemens chunks (3×), a cut from a Siemens rod and two different Si-powders.
Fig. 5. Light microscope images of polished and Secco-etched vertical cuts in the seed-level region of different HPM Si ingots: Reference grown on SIE chunks (top), SIE rod (middle) and 0.3 µm Si-powder (bottom). 118
Journal of Crystal Growth 514 (2019) 114–123
M. Trempa, et al.
Fig. 6. Vertical cut out of the center of a G1 ingot with artificially induced grain boundaries: (a) grain boundary types, (b) crystal orientation parallel to growth direction and (c) growth interface shape measured by lateral photovoltage scanning (LPS). The green arrows indicate grain boundary growth perpendicular to the growth interface, the red arrows a non-perpendicular growth.
energy minimization causing a roof-like structure above the seed stripes [15]. The fact that one roof side is labeled by the Laue-Scanner software as a R grain boundary is according to the used Brandon-criterion [16] justifying the classification of the grain boundaries, although it has to be a Σ3 or Σ9 type due to the crystallographic relationship of the grains in this region. These non-perfect twin boundaries slightly exhibit the tolerable angle deviations larger than 8.7° (in the case of Σ3) and 5° (in the case of Σ9) which results in the labeling as a R grain boundary. Nevertheless, all new formed twin boundaries grow straightly upwards according to the direction of their twin plane and not according to the shape of the solid liquid interface, as can be seen in Fig. 6c and is expected for low index Σ grain boundaries. On the right half of the ingot, seven parallel R grain boundaries were formed by the seeding method. It is obvious that at the beginning of solidification they are growing more or less parallel to each other, however a slight tilt between the grain boundary lines is observable with increasing ingot height. This can be well correlated to the growth interface shape in this ingot region meaning the R grain boundaries propagate perpendicularly to the slightly curved solid-liquid interface as expected (green arrows in Fig. 6c). At about half of the ingot height, two pairs of neighboring R grain boundaries abruptly align each other and coincide with the result that the number of grain boundaries is reduced from four random grain boundaries to two small angle grain boundaries which are not visible in this image. It has to be noticed that in most cases in real HPM growth experiment the annihilation reaction would be R + R → R due to the generally non-equal grain orientation of the involved, next-neighboring grains. This annihilation phenomenon cannot be explained only by a direct correlation to the almost flat shape of the solid-liquid interface, as can be seen by the red arrows in Fig. 6c. In this case the growth direction of the grain boundaries deviates in some degrees from the perpendicular direction to the growth interface, indicating that the
size or increase the random grain boundary fraction at 5 mm ingot height and above in comparison to the reference HPM ingots by the seeding variants presented above. 3.2. Series II: Influence of seeding arrangement on the evolution of R grain boundaries during growth The aim of the experiment series II was to investigate if it is possible to keep alive R grain boundaries in HPM material over ingot height if they were arranged parallel to each other and perpendicular to the crucible bottom, respectively. For that purpose alternating 〈1 1 0〉 and 〈4 2 1〉 oriented seeds were used (arrangement shown in Fig. 2). Because it is well known that the growth interface shape can strongly influence the growth direction of non-twin grain boundaries [17], the growth interface should be as flat as possible. The illustration of the measured growth interface in Fig. 6 c shows an almost flat interface with a slight W-shaped character in the upper part of the ingot. Unfortunately, the exact shape of the growth interface could not be detected in the lower ingot part due to the electrical properties of the silicon seeds. Because the interface becomes more flat towards the top, it can be assumed that the W-shape in the lower part is little bit more pronounced. A vertical cut from the center of this special seeded ingot was investigated by the Laue method. The resulting grain boundary and grain orientation map are shown in Fig. 6a/b. The 10 mm wide seed stripes used for the grain boundary engineering are visible below the seeding level at the bottom of these maps. The ingot can be divided in two parts: In the left half, Σ3 twin boundaries were intentionally induced whereas in the right half, R grain boundaries were generated. Just above the seeding interface, the initially formed Σ3 twin boundaries split into new straightly aligned twin boundaries (examples are marked with white circles), due to 119
Journal of Crystal Growth 514 (2019) 114–123
M. Trempa, et al.
Fig. 7. Top view of seed arrangements and PL-images of vertical cuts of the two G1 ingots grown with thin 〈1 0 0〉 (left) and 〈1 1 1〉 (right) seed stripes. The PLimages are made up of different PL-measurements in order to make the grain boundaries visible over the complete ingot height.
overgrowth of the stripe between the two grain boundaries is triggered by the surface energy of the adjacent grains. In order to validate this hypothesis of competing crystal regions, a further G1 experiment was carried out for which the 〈1 0 0〉 stripes were replaced by 〈1 1 1〉 oriented stripes with 2 mm thickness (see Fig. 7 on the right). The {1 1 1} surface of silicon is known to be the surface with the lowest surface energy. Indeed, the 〈1 1 1〉 oriented stripe region becomes wider with increasing ingot height meaning that the Σ9 grain boundaries are further separated. Interestingly, at ca. 50 mm growth height, the right grain boundary exhibits a kink towards the left grain boundary (marked by the lower red arrow). Laue measurements reveal that this is caused by a twinning process which is well known from the Quasimono technology using monocrystalline seed plates [15]. In this case, the orientation relationship between the 〈1 1 0〉 seeds and the 〈1 1 1〉 seed stripe promotes the twin formation at the edge of the grown 〈1 1 1〉 grain which results in formation of new (twin) boundaries. The initially formed Σ9 grain boundaries are replaced by two Σ3 grain boundaries growing further in upward direction. At approx.70 mm in ingot height the twinning process reoccurs. Again two new Σ3 grain boundaries are formed growing until the top of the ingot. In conclusion, the evolution of the random grain boundaries and all other non-twin boundaries is mainly influenced by the shape of the
kinetics of the surface energy of the adjacent grains plays a more crucial role [18]. Further, it is well known that grains overgrow others with higher surface energy, provided that no strong other influences like highly curved growth interfaces or high growth rates are existent [19]. In the present case, the 〈1 1 0〉 orientation would then have a lower surface energy than the 〈4 2 1〉 orientation (compare grain orientation map in Fig. 6b). Additionally, it can be remarked that the meeting of the grain boundaries in the shown 2D cut is a consequence of other (similar) grain boundary interactions which takes place at lower ingot heights, but at other radial position. A similar observation was made for the other G1 experiment of series II with thin 〈1 0 0〉 oriented seed stripes positioned between 〈1 1 0〉 oriented seed plates (see Fig. 7 on the left). On PL-images of a vertical cut through the ingot center it is visible that both of the left pair of generated Σ29 grain boundaries (type was confirmed by Laue measurements) are grown roughly parallel to each other in the lower part of the ingot up to a certain point at which they were annihilated by overgrowth of the center grain and formation of a small angle grain boundary (SA-GB). On the right, the setup was identical, but the distance between the two grain boundaries was twice as large, i.e. 4 mm instead of 2 mm. This resulted in longer Σ29 grain boundaries which are however also annihilated by overgrowth. The corresponding positions are marked with yellow circles. Also here it is probable that the 120
Journal of Crystal Growth 514 (2019) 114–123
M. Trempa, et al.
h), the growth rate is almost identical at below 5 mm/h at the early beginning of growth which is caused by the thermal inertia of the furnace setup. For the high gas flow rates, differences are already present at the beginning where the growth rate is increased to 9 mm/h and 13 mm/h for 30 m3/h and 50 m3/h gas flow, respectively. The experiment with the cooling gas flow of 10 m3/h equals to the process applied for the growth of a 710 mm high ingot which was simulated by the subsequent growth of eight G1 ingots published in our previous work [4]. For the first ingot, the HPM grain structure was induced by a 20 mm seeding layer containing silicon chunks. For the following seven ingots, a 20 mm thick seed plate cut from the upper part of the previously grown ingot was used. In Fig. 9 the grain boundary length fraction over the total ingot height of 710 mm is shown for the reference experiment (cooling 10 m3/h, growth rate ∼2–10 mm/h) by the squared unfilled symbols [4]. It can be observed that the R grain boundary fraction strongly decreases from 73% to 50% along the first 150 mm grown ingot height and remains constant up to 710 mm, which was interpreted as stable state with annihilation and formation processes being in equilibrium [4]. It has to be noticed that this found equilibrium state is probably linked to the appearing growth conditions in our G1 furnace. In the case of industrial setups the exact values of the grain boundary plateau and the ingot height at which the equilibrium is reached can be different. In the present work we investigated whether a different growth rate would change the grain boundary type distribution over ingot height, especially along the first 150 mm ingot height and within the plateau at ingot heights over 400 mm. For testing the effect of the decreased growth rate of 2–6 mm/h in the lower ingot region, one G1 ingot was grown using silicon chunks as seeds, i.e. starting at an initial ingot height of 0 mm. The effect in the plateau region at this growth rate was investigated by first growing an ingot with a 30 mm thick slice cut from the reference ingot at 370–400 mm height as seed plate, yielding data points at 400 mm and 460 mm. The top of this ingot was then again used as seed for growth of another ingot, allowing to obtain data for a virtual ingot height of 560 mm. The results in Fig. 9 (filled triangle symbols) clearly show that the lower growth rate has no measurable influence on the R grain boundary fraction, both in the lower ingot region as well as in the stable state region. The effect of the higher growth rate of 9–18 mm/h was investigated in a similar way. In this case, three subsequent ingots were grown with the first starting from the silicon feedstock, i.e. from an ingot height of 0 mm. Additionally, one ingot was grown within the stable state region by using a seed plate cut from the reference ingot at 460–490 mm grown ingot height. The results are also plotted in Fig. 9 (filled circle symbols). Within the measurement accuracy again no clear difference to the reference is observable. However, in the first 240 mm ingot height, for all data points a slight decrease of the R grain boundary fraction and a slight increase of twins of 1–3% relative to the trend line of the reference crystal is visible. This trend would agree to the observations of Wong et al. [8] and Lin et al. [9] who proposed an enhanced twin formation with increasing growth rate due to the higher probability of undercooling. Finally, one single ingot was grown with the extremely high growth rate of 12–30 mm/h starting from the silicon feedstock seeding layer (ingot height = 0 mm). Due to the highly concave curvature of the growth interface in the upper parts of the ingot, only two wafers at 35 and 45 mm ingot height could be evaluated. From these results (filled rhombic symbols), again no clear influence of the growth rate on the R grain boundary fraction was observed. Comparably to the other experiment with increased growth rate, the grain boundary type distribution could be slightly shifted to lower R and higher twin values, but again with deviations only within the margin of error. In summary, the realized variation in the maximum growth rate by a factor of 0.5 up to 3 has no significant influence on the grain boundary type distribution. At most, only a slight, but insignificant decrease of the R grain boundary fraction was observed, which we
solid-liquid interface. Additionally, the surface energy of the adjacent grains slightly influences the growth angle of the grain boundaries due to kinetical reasons. The first point is clearly visible also by the “grain boundary arrangement” within the SIE rod experiment in series I. As shown in Fig. 5, the grain boundaries within the rod are arranged in directions from diagonal to parallel with respect to the seeding interface. Despite of this structure, the grain boundaries of the newly formed grains develop more or less perpendicularly to the even flat seeding interface which means their evolution follows the growth interface shape. Based on the above described the following rules to reduce/avoid the annihilation processes of initially formed random grain boundaries can be laid down:
• The initial random grain boundaries should be arranged parallel to each other (perpendicular to the seeding interface). • The distance between neighboring random grain boundaries • • •
shouldn’t be too small in order to shift the meeting of them to larger ingot heights. The growth interface shape should be as flat as possible. The orientations of the adjacent grains should be similar regarding their surface energy. The orientation relationship between adjacent grains should not promote twinning processes.
All these points, with exception of the flat growth interface, are very difficult to achieve within an industrial growth setup. In particular, some of the mentioned points are in contradiction to the use of randomly oriented seeds as it is applied for the HPM process. Even though artificially arranged seeding layers yielding a high and stable random grain boundary fraction might exist, the required seed arrangement would be even more complex than for the Quasi-Mono approach. Therefore, this approach is probably no realistic option from an economic point of view. 3.3. Series III: Influence of growth rate on the growth behavior of R grain boundaries To investigate the influence of the growth rate, four G1 ingots were crystallized with argon cooling gas flow rates of 0 m3/h, 10 m3/h, 30 m3/h and 50 m3/h. The resulting growth rates for these conditions are depicted in Fig. 8. As intended, an increasing gas flow results in a higher growth rate. In the case of low gas flow rates (0 m3/h and 10 m3/
Growth rate [mm/h]
40
50 m³/h 30 m³/h 10 m³/h 0 m³/h
30
20
10
0 0
20
40
60
80
100
120
Ingot height [mm] Fig. 8. Measured growth rate over the grown ingot height for varying cooling gas flow conditions. 121
Journal of Crystal Growth 514 (2019) 114–123
grain boundary length fraction [%]
M. Trempa, et al.
80
R Σ3
70 60
2-6 mm/h
50
2-10 mm/h (Ref.)
40
9-18 mm/h
30
12-30 mm/h
20 10 0
0
50
100
200 300 400 500 600 700
grown ingot height [mm] Fig. 9. Grain boundary length fraction over the grown ingot height for different growth rates.
growth rate of 2–6 mm/h at the ingot heights 460 mm and 560 mm. This difference is also obvious to the naked eye from the PL images on the right hand side in Fig. 10. Since the initial dislocated area in these ingots should be about 12% (the seed plate was cut at 370–400 mm from the reference ingot), the dislocated area is decreased down to 8%. This indicates that some of the existing dislocation clusters were overgrown by dislocation free grains. Maybe also less thermal stress is generated by the lower growth rate, which reduces the dislocation multiplication.
assess as deterioration instead of an improvement relative to the reference process. Even though this might indicate that even lower growth rates than the 2–6 mm/h tested here could yield a (slightly) increased R grain boundary fraction, at least in the lower region of the ingot, such an approach would not be feasible in industry due to economic considerations. Thus, the authors suggest that the process parameter growth rate is no tool to further improve the random grain boundary fraction in industrially produced HPM silicon ingots. Even though the grain structure has been proven to be a very good indicator for the resulting solar cell performance, it is only an indirect method relying on the correlation between grain structure and emergence of dislocations. Therefore, beside the grain structure analysis, the wafers were also investigated by means of PL imaging to determine the recombination active dislocation clusters. In Fig. 10, the recombination active wafer area, occupied by sum-up the area covered by dislocation clusters (grey-value analysis) divided by the total wafer area, is plotted along the grown ingot height. For the highest growth rate of 12–30 mm/h no statement can be made due to the lack of samples in ingot heights > 45 mm. Further, an increased growth rate of 9–18 mm/h yields a similar trend of the recombination active area over the ingot height as the reference. In contrast, a reduced dislocated area relative to the reference is measured for the lower
4. Conclusion Several hypothetical approaches to improve the grain structure properties of HPM silicon ingots were evaluated by experiments in G1 scale. It was shown that the use of extremely fine Si particles or densely packed chunk layers can lead to a finer initial grain structure and probably also to a higher length fraction of the suitable random grain boundaries, but only in the first few 100 µm of growth. The natural grain selection results in a quick grain coarsening and therefore no positive effect remains at relevant block heights. From grain boundary engineering experiments we found some directives, like a parallel arrangement of the initial random grain
Fig. 10. Recombination active area over grown ingot height for different process conditions. On the right hand side the PL-images of three ingots at ∼550 mm ingot height are shown. 122
Journal of Crystal Growth 514 (2019) 114–123
M. Trempa, et al.
boundaries or keeping flat the growth interface, which may help to reduce the annihilation of the random grain boundaries. However, an industrial implementation seems to be rather complicated, especially concerning the even complex seed arrangement which has to fulfill a lot of requirements. Finally, also the growth rate, the influence of which was tested over a wide range, has no influence on the annihilation processes of the random grain boundaries. Nevertheless, the area fraction of electrically active dislocations was found to be reduced for applying a reduced growth rate, which again is unfortunately undesirable in industry processes for economic reasons. In summary, it has to be recognized that from the grain structure point of view, the properties of the current HPM process are most probably the best ones which can be achieved, at least under industrial conditions, and therefore further improvements of HPM material quality has to be approached via other paths, e.g. the reduction of thermal stress or of impurities like metals or oxygen.
References [1] ITRPV, 2018. International Technology Roadmap for Photovoltaic (ITRPV), 2017. ninth ed. [2] C.W. Lan, W.C. Lan, T.F. Lee, A. Yu, Y.M. Yang, W.C. Hsu, B. Hsu, A. Yang, J. Crys. Growth 360 (2012) 68–75. [3] Y.M. Yang, A. Yu, B. Hsu, W.C. Hsu, A. Yang, C.W. Lan, Prog. Photovoltaics. Res. Appl. 23 (3) (2015) 340–351. [4] M. Trempa, I. Kupka, C. Kranert, T. Lehmann, C. Reimann, J. Friedrich, J. Crys. Growth 459 (2017) 67–75. [5] R.R. Prakash, K. Jiptner, J. Chen, Y. Miyamura, H. Harada, T. Sekiguchi, Appl. Phys. Express 8 (3) (2015) 35502. [6] C. Reimann, M. Trempa, T. Lehmann, K. Rosshirt, J. Stenzenberger, J. Friedrich, K. Hesse, E. Dornberger, J. Crys. Growth 434 (2016) 88–95. [7] K.E. Ekstrøm, G. Stokkan, A. Autruffe, R. Søndenå, H. Dalaker, L. Arnberg, M. Di Sabatino, J. Crys. Growth 441 (2016) 95–100. [8] Y.T. Wong, C. Hsu, C.W. Lan, J. Crys. Growth 387 (2014) 10–15. [9] H.K. Lin, M.C. Wu, C.C. Chen, C.W. Lan, J. Crys. Growth 439 (2016) 40–46. [10] W. Chen, Q. Wang, D. Yang, L.D. Li, X.G. Yu, L. Wang, H. Jin, J. Crys. Growth 467 (2017) 65–70. [11] M. Beier, M. Trempa, J. Seebeck, C. Reimann, P. Wellmann, B. Gründig-Wendrock, J. Friedrich, K. Dadzis, L. Sylla, T. Richter, in: Proc. of the 29th European Photovoltaic Solar Energy Conference, Amsterdam, pp. 717–721. [12] C. Huang, H. Zhang, S. Yuan, Y. Wu, X. Zhang, D. You, L. Wang, X. Yu, Y. Wan, D. Yang, Sol. Energ. Mat. Sol. C. 179 (2018) 312–318. [13] T. Lehmann, M. Trempa, E. Meissner, M. Zschorsch, C. Reimann, J. Friedrich, Acta Mater. 69 (2014) 1–8. [14] M. Trempa, C. Reimann, J. Friedrich, G. Müller, A. Krause, L. Sylla, T. Richter, Cryst. Res. Technol. 50 (2015) 124–132. [15] M. Trempa, C. Reimann, J. Friedrich, G. Müller, D. Oriwol, J. Crys. Growth 351 (2012) 131–140. [16] D.G. Brandon, Acta Metall. 14 (11) (1966) 1479–1484. [17] Klaus-Thomas Wilke, Joachim Bohm (Hg.), Kristallzüchtung. 1. Aufl. [18] T. Duffar, A. Nadri, Crystal Growth/Croissance Cristalline 14 (2) (2013) 185–191. [19] K. Fujiwara, Y. Obinata, T. Ujihara, N. Usami, G. Sazaki, K. Nakakjima, J. Crys. Growth 266 (2004) 441–448.
Acknowledgement These research and development activities are partly carried out within the ENOWA-II project (0325805E) funded by the German Federal Ministry for Economic Affairs and Energy. Appendix A. Supplementary material Supplementary data to this article can be found online at https:// doi.org/10.1016/j.jcrysgro.2019.03.005.
123