Stress and doping uniformity of laser crystallized amorphous silicon inthin film silicon solar cells

R. M. B. Agaiby,1,a� M. Becker,1,2 S. B. Thapa,3 U. Urmoneit,4 A. Berger,1,2 A. Gawlik,2

G. Sarau,2 and S. H. Christiansen1,2

1Max-Planck Institute of Microstructure Physics, Weinberg 2, 06120 Halle, Germany2Institute of Photonic Technology, Albert-Einstein-Str. 9, 07745 Jena, Germany3Max-Planck Institute for the Science of Light, Günther-Scharowsky-Str. 1, 91058 Erlangen, Germany4Bavarian Laser Center GmbH, Konrad-Zuse-Straße 2-6, 91052 Erlangen, Germany

�Received 6 November 2009; accepted 19 January 2010; published online 8 March 2010�

Simultaneous and locally resolved determination of the mechanical stress variation and the free holeconcentration using Raman spectroscopy is demonstrated in laser crystallized amorphous siliconlayers. Such layers are often used for the fabrication of thin film solar cells, e.g., on borosilicateglass substrates. The combined effects of stress and doping on the Raman signal can be separatedbased on the use of three wavelengths in the visible. The results show that the free holeconcentration in the samples investigated varies between 1�1018 and 1.3�1019 cm−3. Stress aswell as the free hole concentration vary substantially within the sample. The stress level variesbetween 575 and 850 MPa ��12 MPa�. Cross-sectional transmission electron microscopy imagesshow the presence of extended lattice defects such as dislocations and grain boundaries in thecrystallized Si layer which could account for the lateral stress variations detected by Ramanspectroscopy. The impact of film inhomogeneity in terms of stress and doping on the performanceof a solar cell will be discussed. © 2010 American Institute of Physics. �doi:10.1063/1.3319654�

I. INTRODUCTION

Thin film solar cells �TFSC� are quickly becoming apopular alternative to the more costly crystalline silicon solarcells. While the latter experience better efficiency scales incomparison with TFSC, their cost does not justify continuedhigh scale production. Therefore, optimizing TFSC to pro-duce comparable efficiencies has been one of the main aimsof the photovoltaics industry. Low cost solar cells based onamorphous silicon �a-Si� have not yet been able to performin terms of efficiency as well as crystalline silicon wafersolar cells.1 Therefore, a good compromise could be to crys-tallize a-Si thin films, e.g., on low-cost substrates like glass.This procedure has the advantage of using less of the highquality, low impurity silicon and, therefore, costs less. TFSCon glass are fabricated by starting with an a-Si seed layer thatcan easily and homogeneously be deposited by vapor phasedeposition methods such as chemical or physical vapordeposition.2 The a-Si is then crystallized in a second process-ing step to finally obtain a multicrystalline Si seed layer tem-plate for further solar cell processing. Crystallization of a-Si�preferably doped to a level that is needed according to theapplied solar cell concept� can be carried out, e.g., by a fur-nace anneal3 or by using a pulsed4 or continuous wave �cw�laser5 to supply the energy needed to take the a-Si layer fromits metastable amorphous state to the stable crystalline state.In this paper we will report on cw-laser crystallized microc-rystalline Si seed layers on glass closely resembling the ma-terial reported in.6 The material is characterized by grains ofdifferent sizes depending on which region in the crystallizedline, analysis is carried out. The center of the line shows

larger grains than at the rims due to a diode laser energyprofile within the laser focus. The material shows varyinginternal stresses and dopant concentrations depending on thelocal microstructure. The role of the laser crystallized layerfor the TFSC is twofold: �1� it serves as a seed layer forfurther epitaxial thickening to form an entire solar cell com-posed of absorber, emitter, and contact layers all to be inte-grated on the glass substrate and all to be composed of mul-ticrystalline Si �Ref. 6�; �2� moreover, it serves as a highlydoped contact layer that has metal like conductivities due todegenerate doping. Given these roles it is important to ac-count ideally for the largest possible grain size with the leastpossible internal stresses and extended lattice defect densi-ties, in addition to the highest possible homogeneity of dop-ants throughout the highly doped, multicrystalline seed layer.Thus, it is important to know the overall stress level that isincorporated in the laser crystallized layer. Stress is deter-mined by the temperature dependent thermal mismatch be-tween Si and the glass substrate3 and by details of the ther-mal budget involved in the laser processing. Stress is alsoaffected by the lattice defect population that is responsiblefor stress release or stress increase being locally superim-posed by the overall internal stress induced by the thermalmismatch.7 Internal stresses vary throughout the layer, inter-acting with lattice defects and affecting stress induced dopantdiffusion and thus dopant inhomogeneity.8 Quantifying thestress at or near different types of grain boundaries within theseed layer and correlating it with the dopant distribution per-mits optimizing the thin film material needed to fabricatehigh quality multicrystalline TFSC.

Secondary ion mass spectroscopy has been used as anestablished method in determining doping profiles in silicondue to its high sensitivity and reproducibility; however, it isa�Electronic mail: agaiby@mpi-halle.mpg.de.

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destructive and time consuming and lacks good spatial reso-lution which is in the several square micrometer range atbest.9 Alternatively, x-ray diffraction is capable of determin-ing stresses and doping in materials,10 however, it also lacksspatial resolution and can be quite challenging to be appliedto thin films.11 Raman spectroscopy is a potential contenderin simultaneously determining stresses and doping concen-trations due to its high spatial resolution, repeatability, andnondestructive nature.12,13 The spatial resolution of Ramanspectroscopy can be improved further, from the micrometerto the nanometer scale, by bevelling samples at very shallowangles ��0.5°� thereby acquiring high spatial resolution pro-files at least in the direction of the wedge.14

Previous studies have shown that heavy p-type doping,e.g., boron, in silicon results in Fano resonances which arecharacterized by the shifting of the Raman Si peak to lowerwavenumbers as well as by a peak asymmetry on the higherwavenumber side of the peak.15 For heavy n-type doping,e.g., by phosphorous, the peak asymmetry resides on thelower wavenumber side of the peak.16 Peak asymmetry hasbeen attributed to the overlap of the transition probabilitiesbetween continuum �electronic� and discrete �phonon�scattering.17 The extent of peak asymmetry defined by thesymmetry parameter q has been used by Cerdeira et al.15,18

and others16,19,20 to directly determine the free carrier con-centrations. However, there is a large discrepancy in litera-ture between the symmetry parameters q stated for differentdopant concentrations. Therefore, a calibration of the relationbetween the q parameters and the free carrier concentrationsmay need to be repeated.

First order Raman silicon peaks are largely influenced bygrain orientation,21 grain size,22 stress,12 and doping.15,23 Si-multaneously quantifying stress, doping, grain size, and ori-entation is therefore not trivial. Silicon peak asymmetrycharacterized by the symmetry parameter q is determined byfitting the peak to the equation describing the Fanoresonances15 thereby giving a direct indication of the freecarrier concentration. Peak shift due to stress can then bedetermined by eliminating the shift due to doping from theoverall peak shift. Grain orientation can be determined bythe difference in silicon peak intensity, acquired under vari-able polarization settings.24

This work attempts to thoroughly investigate laser crys-tallized a-Si layers on glass, bevelled at very shallow angles��0.2°�, using Raman spectroscopy in order to isolate thevarious influences of doping and stresses from the character-istic silicon peaks, thereby precisely quantifying each contri-bution. It examines material homogeneity in terms of theuniformity of the doping concentration and the stress value,in order to detect potential performance deterioration ofTFSC fabricated by laser mediated a-Si crystallization.

II. EXPERIMENTAL DETAILS

Two samples were used in this study; �1� a calibrationsample to determine the relation between free carrier/dopantconcentration and the symmetry parameter q and �2� the seedlayer of a TFSC which is analyzed with respect to stress

distributions and carrier concentration fluctuations within thecrystallized a-Si layer depending on the local lattice defectstructure/distribution/population.

The calibration sample was grown by depositing�400 nm of a-Si with about 4�1019 cm−3 boron doping ona borosilicate glass substrate, Borofloat from Schott AG,25

using electron beam evaporation at 200 °C. The a-Si layerwas crystallized with a LIMO line focus diode laser �linelength: 1 cm at 130 �m line width; wavelength: 808 nm;maximum power: 320 W� that shows a much more homoge-neous energy profile especially along the crystallizing laserline6 �discussed in Sec. III A�. The calibration sample is usedto correlate the symmetry parameter q to the respective firstorder Raman peak shift in silicon and to the free hole con-centration. The sample was bevelled in order to expose areaswith different doping thereby sufficiently spreading out thethin multicrystalline layer, exposing the interfacial area rightnext to the glass. It also accommodates for enough measure-ment points by Raman spectroscopy within each dopinglevel. The bevelling procedure is described below in thisexperimental section and in Fig. 1. The seed layer of theTFSC was grown by depositing about 532 nm a-Si layer withabout 4�1019 cm−3 boron doping and was crystallized us-ing a cw-diode laser26 operating at a wavelength of 808 nmand a power of 4.2 W. The calibration sample and the ana-lyzed sample were placed on a hot stage and heated to atemperature of 650 °C and 610 °C, respectively, to avoidlayer or substrate cracking upon laser melting of a-Si andrapid resolidification. The scanning speed for the calibrationsample and the analyzed sample was 3 cm/s and 25 mm/s,respectively, in ambient air during crystallization. Scanningwas carried out with a high precision positioning stage. Theapplied laser power �given that the wavelength is fixed andthe laser focus resides ideally on the sample surface and nocap or interfacial layers were used between the glass and thea-Si or on the a-Si layer�, the temperature of the hot plate aswell as the scan speed determine the grain size.27 These pa-rameters were chosen to prevent thin film damage at veryhigh laser power since the diode laser crystallization processis unstable, the thinner the a-Si layer is. The instabilitiesoccur due to two reasons. First, there is a tendency of dew-etting when silicon is liquid. Dewetting is faster in thinnerlayers and leads to holes in the crystallized layer. Second, theabsorption of a-Si, at least at room temperature, is low forthe diode laser wavelength used so that only a part of thelaser energy is effectively used to heat and eventually meltthe a-Si.

FIG. 1. �Color online� Schematic of the beveled sample showing the ex-tended film thickness �stretching factor� of the thin crystallized a-Si layer asa result of beveling. Raman measurements taken along the wedge increasethe resolution of Raman spectroscopy to the nanometer scale.

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For the bevelling, the samples �calibration sample andseed layer of the TFSC� were placed on a wedge in a resinmold and bevelled using a Buehler polisher, mechanicallythen chemically for a smooth polish. The samples were bev-elled in order to expose very thin layers over larger areasthereby, improving the depth resolution of Raman spectros-copy �Fig. 1�. The bevel angle � of the seed layer of theTFSC was less than 0.2° resulting in a stretching factor S���of about 320 �formula given in Fig. 1�. After bevelling, thesamples were etched in diluted Secco etch28 for 15 s to op-tically reveal grain boundaries and defects.

Raman measurements were taken with a Horiba JobinYvon LabRam system equipped with an Ar+ ion laser �wave-lengths: 488 and 514 nm� and a HeNe laser �633 nm�. Theselasers were used for generating calibration curves correlatingthe symmetry parameter q with the silicon peak shift and thefree carrier concentration. For the analyzed multicrystallineSi layers, the three lasers were used for acquiring maps of thesame region to determine the distribution of stress and freehole concentration. The Raman curves of the same area ap-pear differently taken with excitation at different wave-lengths �Fano resonances are known to diminish with de-creasing excitation laser wavelength29 which means that thecorresponding peak asymmetry becomes less pronounced� sothat measurements with three different wavelengths result inthree independent experiments. The penetration depth for the633, 514, and 488 nm lasers are 2–3 �m, 770 nm, and 570nm, respectively.12 The Raman system was calibrated beforeeach map acquisition in order to prevent accumulated timedependent machine drift. Low laser power ��6 mW� wasused in order to prevent localized sample heating which canresult in significant peak shifts.30–33 The lateral step size ofthe map acquisition in the x-y direction was 2.5 �m result-ing in a projected depth resolution of 8–10 nm due to theshallow angle bevelling �Fig. 1�. The error in peak positiondue to fitting is about �0.05 cm−1 �recorded for differentfitting iterations� resulting in an equivalent error in stress of�12.5 MPa.

III. RESULTS AND DISCUSSION

A. Calibration sample

Figure 2�a� shows an optical image of the bevelled andlaser crystallized seed layer ��400 nm in thickness and 4�1019 cm−3 boron doped�; the layer appears to be homoge-neous at the displayed scale, i.e., no grain boundaries or de-fects are visible even after short defect etching with Seccoetch.28 Figure 2�b� is a silicon peak asymmetry mapping,acquired with the 633 nm laser, depicting the distribution ofthe free hole concentration in the bevelled seed layer of thecalibration sample. Lighter regions experience higher peakasymmetry, on the high wavenumber side of the peak, i.e.,according to the boron doping used, in comparison with thedarker regions. This implies that lighter regions have higherfree hole concentrations as opposed to the darker regions,i.e., the nominally homogeneous starting doping level issmeared out from the glass/Si interface toward the samplesurface. The smearing out occurs during laser liquifying ofthe a-Si and its recrystallization. The homogeneous doping

profile lines perpendicular to the wedge, parallel to the focusline show that, laterally, the energy intake and doping redis-tribution upon crystallization was very homogeneous. Due tobevelling the initially �400 nm thick layer appears to beabout �40 �m thick. The symmetry parameter q in lines ofidentical color �i.e., identical asymmetry of the peak and thusfree hole concentration� was extracted by fitting the siliconpeak to the equation describing the Fano resonances which isgiven by18

I =

I0 � ��q + � 2

� � �� − �MAX�2�

�1 + � 4

�2 � �� − �MAX�2+ b �1�

where I is the peak intensity �a.u.� along the y-axis, � is thewavenumber �cm−1� along the x-axis, �MAX is the siliconpeak position under the influence of doping, I0 is a scalingfactor, b is baseline correction, and � is the peak width�cm−1�. Figure 3 shows the silicon peaks extracted at four

FIG. 2. �Color online� �a� Optical image of the beveled calibration sample��400 nm crystallized a-Si on borosilicate glass �Ref. 25�� after Seccoetching showing the homogeneity of the sample �no grain boundaries andother defects� and �b� the Raman peak asymmetry map of the beveled cali-bration sample, acquired using the 633 nm laser, showing the variation inthe free hole concentration within the sample depth. Points marked on themapping indicate the areas for which the symmetry parameter q was ex-tracted and correlated with the corresponding silicon peak position and freecarrier concentration �see Figs. 4 and 5�.

FIG. 3. First order Raman Si peaks extracted at different points on thebeveled calibration sample captured with the 633 nm laser. As the dopingincreases, the peak asymmetry increases �Fano resonances� while the peakintensity decreases.

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different points on the sample �marked in Fig. 2�b��. As ex-pected, the silicon peaks become more asymmetric with anincrease in doping and the peak intensity decreases accord-ingly. Peak asymmetry maps similar to that in Fig. 2�b� werealso acquired with the 488 and 514 nm lasers. Figure 4�a�shows the correlation between the symmetry parameter q andthe respective silicon peak positions extracted from the threemaps taken at different excitation wavelengths at the samearea. This peak shift is considered to be due to doping only,i.e., the sample is assumed to be stress free due to the samplehomogeneity which is observed in Fig. 2�a�. During the crys-tallization process, the sample stage is set to about 610 °C sothat the borosilicate glass softens to accommodate thestresses incorporated in the crystallizing a-Si layer and thesubstrate �glass� during the melting and cooling process. Ex-trinsic stress which could be introduced to the sample as aresult of any thermal mismatch between the borosilicateglass and the crystallized a-Si is about 3 MPa �Ref. 34�which falls within the experimental error of the Raman mea-surements.

The expected theoretical peak shift due to doping wasalso determined using the experimentally fitted parametersand the following equation18

�MAX = �0 +��

2q�2�

where �0 is the peak position for bulk undoped and stressfree silicon �cm−1� and �� is the change in peak width due to

doping �cm−1�. Comparison between Figs. 4�a� and 4�b�shows that the trend in peak shift with respect to q betweenthe experimental and the theoretical curves agree, although,the experimental peak shift due to doping is higher than theestimated theoretical peak shift. This difference could be ex-plained by the approximations that are usually applied tosimplify theoretical calculations. For low values of q �i.e.,high doping�, the silicon peak position decreases, however, itsaturates with increasing q �i.e., low doping�. Fano reso-nances are known to diminish with decreasing excitation la-ser wavelength29 which means that the corresponding peakasymmetry becomes less pronounced. This is because thedependency of the electronic and phonon scattering on theexcitation laser wavelength is different, thereby causing theratio of the interaction between the scatterings to changewith wavelength.29 Therefore, the correlation between q andthe respective silicon peak positions acquired using the dif-ferent lasers is not the same.

Figures 5�a�–5�c� correlate the q parameters acquiredwith the 488, 514, and 633 nm lasers, respectively, at thepoints shown in Fig. 2�b�, to the free hole concentration dueto boron doping. Values for the free hole concentration wereobtained from18,19 by taking into account the common values

FIG. 4. �a� Experimental and �b� theoretical calibration curves correlatingthe peak shifts due to boron doping in silicon to the symmetry parameter qrepresenting peak asymmetry. Trends between theory and experiment agree.However, there is a slightly larger change in the peak shift exhibited in theexperimental data.

FIG. 5. Free hole concentration correlated with the inverse of the symmetryparameter q extracted at different depths in the calibration sample capturedwith the �a� 488 nm, �b� 514 nm, and �c� 633 nm lasers. Free hole concen-tration was determined from. �Refs. 18 and 19�

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of q between literature and the current work. As can be seen,the trend line shows a linear dependence between the freehole concentration and the inverse symmetry parameter q forall the three laser wavelengths. The slope for the 488 nmlaser is the steepest followed by the one for the 514 and 633nm lasers. Doping in the calibration sample decreases gradu-ally toward the sample surface and varies from 6.5�1018 cm−3 �point 8� to 5�1019 cm−3 �point 1�. These cali-bration curves can then be used to estimate the free holeconcentration using the values of q acquired from othersamples with any of the three lasers. As an example of tech-nological interest for future TFSC concepts, the cw-diodelaser crystallized silicon seed layer of a TFSC was analyzed.

B. TFSC seed layer

Figure 6 shows optical images of the bevelled and Seccoetched laser crystallized a-Si seed layer. As can be seen inFig. 6�a�, grains of various sizes form with larger grains re-siding in the laser crystallized line center where the energytends to be maximum and smaller grains forming at the linerims where energy of the crystallizing laser beam is smalleraccording to the laser beam Gaussian energy profile. In theline center the grain size varies in the range of 10–100 �mdepending on growth anisotropy of grains of different orien-tation. The overall size of the large center grains is deter-mined by the highest energy density there and thus the low-est cooling rate as well as the lowest nucleation rate.5

Consequently, low heat density at the rim of the laser lineresults in faster cooling rate and therefore smaller grains�Fig. 6�b��. Since most of the grain boundaries are recombi-nation centers for carriers that are created by photons of suf-ficiently high energy35 these extended defects deteriorate thesolar cell performance and therefore should be avoided.Therefore, during laser crystallization, a sufficiently large

overlap between laser crystallized lines needs to be main-tained, in order to avoid the small grained areas. Thus, ide-ally a line focus laser as the one used to crystallize the cali-bration sample rather than a spot focus �with a Gaussianprofile� are ideal to obtain large grained, low defect density,and homogeneously doped material.6 For our Raman experi-ment, however, the test sample with well separated Gaussianspot focus laser lines is a good enough model sample, espe-cially to study the interplay of lattice defects such as grainboundaries with dopant segregation and stress states relatedto such extended lattice defects. Figure 6�c� shows the grainstructure in our model sample, where the crystallized stripesare about 50 �m apart thus showing there was no overlapbetween successive laser line scans.

Figure 7�a� shows the optical image of the Ramanmapped region that was analyzed for doping and stress. Fig-ure 7�b� shows the peak intensity map captured with the 514nm laser. Similar maps of the same region were acquiredusing the 488 and 633 nm lasers. Polarization measurements

FIG. 6. �Color online� Optical images of the bevelled and Secco etched laser crystallized a-Si seed layer showing �a� a large crystallized stripe, �b�microcrystals formed at the edge of the stripes, and �c� the distance between successive well separated laser crystallized stripes showing no overlap betweenindividual laser line scans during crystallization.

FIG. 7. �Color online� �a� Optical image of the Secco etched laser crystal-lized a-Si seed layer and �b� the Raman intensity map acquired using the 514nm laser showing multiple grains with different sizes. Grain boundaries inthe maps correlate well with the optical image. Points on the maps wereused to compare the stress magnitudes acquired using the different lasers.

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were used in order to easily identify grain boundaries withinthe mapped region, since Raman peak intensities depend ongrain orientations next to its dependence on stress and dop-ing levels.21 As can be seen, the Raman peak intensity mapagrees well with the grain boundaries visible in the opticalimage. Points marked on the intensity map refer to the areasthat were analyzed for precise free hole concentration andstress quantification.

1. Free hole concentration

Peak asymmetry maps showing the distribution of thefree hole concentration in the seed layer are shown in Figs.8�a�–8�c� captured with the 488, 514, and 633 nm lasers,respectively. The distributions of the free hole concentrationcaptured with the three lasers agree even though the peak

asymmetry decreases with the excitation laser wavelength.However, there is a slight machine table drift which is de-picted in the maps as an offset in the captured patterns. Theredoes not seem to be any evidence of dopant distribution de-pendency on the grain boundaries. As can be seen in Fig. 8,the free hole concentration in the captured maps varies be-tween 1�1018 and 1.3�1019 cm−3. This range is muchlower than that intended during processing �4�1019 cm−3�.However, the variation in doping within the 35�40 �m2

mapped area by about one order of magnitude shows theseverity of the doping inhomogeneity. This range is repro-ducible in all the three maps thus showing that the measure-ments are reproducible with the three lasers within experi-mental error. Carrier concentration values were extracted bycorrelating the q parameters, extracted from the peak asym-metry maps, to the calibration curves in Fig. 5. Figure 9shows the Raman spectra captured with the 633 nm laser atpoints 9 and 10 �marked in Fig. 7�b�� where there is evidenceof increase in peak asymmetry due to higher doping at point10 �1.3�1019 cm−3� in comparison with point 9 �2�1018 cm−3�. The peak intensities have been normalized toclarify the difference in peak asymmetry.

2. Stress variation

After determining the distribution of the free hole con-centration, ten points were chosen for the evaluation of thestress distribution. Each peak had to be individually fittedusing Eq. �1� in order to determine the contribution of thedoping �from q� to the peak shift using Fig. 4�a�. The remain-ing peak shift was attributed to stress. Assuming biaxial iso-tropic stress, it was calculated using the following equation12

= 250 � ��MPa� �3�

where is the stress and � �cm−1� is the silicon peak shiftdue to the stress. Table I shows the experimental stress val-ues extracted at the ten points in each of the three mappingsacquired using the 488, 514, and 633 nm lasers. As can beseen, the stress values extracted from the maps captured withthe three lasers agree to a large extent. There is a variation ofabout 275 MPa in stress within the mapped region showingevidence of lateral stress gradients across the sample.

FIG. 8. �Color online� Raman peak asymmetry maps taken with �a� 488 nm,�b� 514 nm, and �c� 633 nm lasers, showing the agreement of the distributionof the free hole concentration in the three maps. There is a slight machinetable drift between the three maps depicted by a slight change in the cap-tured patterns.

FIG. 9. Raman spectra captured with the 633 nm laser showing the increasein peak asymmetry with increase in doping between points 9 and 10. Freecarrier concentration is 2�1018 cm−3 at point 9 and 1.3�1019 cm−3 atpoint 10. Points are marked in Fig. 7�b�. Peak intensities have been normal-ized for clarity.

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Stress in the sample occurred mainly due to stress fieldsinduced by extended lattice defects such as dislocations andgrain boundaries.36 The laser crystallization process was con-ducted such that the thin silicon layer on the glass substratecould crystallize at the heating stage temperature at above610 °C, a value for which the borosilicate glass is above itsannealing point so that the glass can accommodate any stressthat may arise due to the difference in the coefficients ofthermal expansion. In the presence of any thermal �extrinsic�stress between the thin film layer and the substrate due to thedifference in the coefficient of thermal expansion betweenthe two,25,37,38 the extrinsic stress can be calculated theoreti-cally using the following equation39

th = �af − as���T�Ef

1 − yp�4�

where af and as are the average coefficient of thermal expan-sion of the film �silicon� and the substrate �glass�, respec-tively. Ef is Young’s modulus of the film �1.69�1011 Pa�, yp

is Poisson’s ratio �0.27, for Si�111��, and �T is the differencebetween room temperature and the temperature at which theglass substrate turns soft/liquid. For borosilicate glass thistemperature is the annealing point at 560 °C �Ref. 25� atwhich the glass can accommodate the thermal stresses pro-duced inside it as well as in the crystalline Si layer. UsingEq. �4�, the theoretical overall, extrinsic stress due to thethermal mismatch between the Si layer and the glass sub-strate is about 3 MPa �Ref. 34� which is within the experi-mental error of the Raman measurements. The stress valuesshown in Table I taken at the ten measurement points indi-cated in Fig. 7�b� show values that range between approxi-mately 580 and 850 MPa ��12 MPa�. These values occurdue to the local grain boundary structure and stress fieldsinduced by extended lattice defects such as dislocations. Thestress values differ when taken at the three different excita-tion wavelengths due to the slight stage drift between thethree measurements resulting in a neighboring region �i.e.,not exact points between the three maps� at which thestresses were extracted.

3. Electrical measurements

In order to analyze the electrical conductivity of the seedlayer, current-voltage �I-V� measurements were carried outusing a Semiconductor Parameter Analyzer HP4156B byforming a direct contact with micromanipulators. The contactdistance between the micromanipulators on the layer waskept fixed at 1 mm for all the measurements. Before per-forming the measurements, the sample �before beveling andSecco etching� was put into HF �2%� dip for about a minuteto remove the native oxide that has formed during the lasercrystallization process. Figure 10�a� shows the optical mi-croscopy image of the laser crystallized seed layer. The num-bers �1 and 2: laser crystallized strips, 3 and 4: edge of thestrips� indicate the area where electrical measurements wereperformed. Figure 10�b� illustrates the two-point I-V curvesmeasured at different areas on the layer �corresponding toFig. 10�a��. The inconsistency of the plots can be ascribed tothe varying grain sizes and the density of grains, presence ofdefects, and the inhomogeneity of the doping distribution atdifferent areas of the layer as deduced from Raman dopingmappings. As expected, a relatively better electrical conduc-tivity was measured at the crystallized stripe where thegrains are large �plot 1�. However, the conductivity de-creased with a decrease in grain size �see plot 2�. This maybe attributed to the increased carrier scattering at grainboundaries that is more pronounced in the small grained ar-eas of the crystallized lines �plot 2�. Similarly, we obtainedeven lower electrical conductivity at the edges of the stripes�plots 3 and 4� where even smaller microcrystals wereformed, i.e., where the grain boundary density was evenhigher. Unfortunately, a more reliable four-point I-V mea-surement could not be carried out due to the very large sur-face roughness of the sample which is a result of the meltmediated crystallization process �a volume increase occursupon crystallization so, that surface bumps form at pointswhere the melt crystallizes the latest�.

4. Cross-sectional transmission electron microscopy„XTEM… investigations

XTEM investigations were carried out to check the ho-mogeneity of the laser crystallized seed layer. As seen in the

TABLE I. Stress values �MPa� extracted from selected points in the map-pings acquired using 488, 514, and 633 nm lasers. Stress values taken withthe three different excitation wavelengths differ due to the slight stage driftbetween the three measurements resulting in a neighboring region �i.e., notexact points between the three maps� at which the stresses were extracted.

488 nm 514 nm 633 nm

1 762.5 587.5 672.52 757.5 747.5 682.53 677.5 600 637.54 847.5 632.5 612.55 772.5 592.5 597.56 510 667.5 6107 670 700 647.58 697.5 675 6259 575 692.5 647.510 762.5 650 682.5

FIG. 10. �Color online� �a� Optical image of the crystallized seed layershowing four points at the rims and centers of the crystallized laser line atwhich two-point I-V curves were captured and �b� I-V curves taken at fourpoints showing lower electrical conductivity at points 3 and 4 in comparisonwith points 1 and 2; this fact can well be attributed to the grain sizesinvolved and the dopant concentration variation at the different measure-ment points. I-V measurements in the amorphous green stripe, between thelaser crystallized lines, are not plotted as the measured current was too lowto be visible on the same axis.

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XTEM image of Fig. 11, the crystallized layer is about 500nm thick with some of the defects extending throughout thelayer, from the glass substrate to the layer surface. Figure 11shows straight grain boundaries that are determined to betwin boundaries, the most commonly present defect in lasercrystallized silicon.6 Other grain boundary types such as ran-dom grain boundaries and coincident site latticeboundaries35,40 are also present in the material and are per-ceived to be more harmful for the solar cell. These grainboundaries have a higher recombination activity of carriersand are often superimposed by additional dislocationnetworks36 that accommodate internal stress and that are usu-ally highly recombination active themselves.41 The circulardefects in Fig. 11 are dislocations/dislocation loops that alsoform to accommodate stresses and that are characterized by astress field themselves.

C. Discussion of experimental results

Using Raman spectroscopy we determined inhomoge-neous free hole concentration distributions and inhomoge-neous stress distributions in cw-diode laser crystallized a-Silayers on glass substrates that are used in novel thin filmsilicon solar cell concepts on glass.6 We find extended latticedefects such as dislocations, dislocation networks, and grainboundaries of various types in the present multicrystallinesilicon thin film material. These lattice defects are found toaffect the local stress states in the layer and the local dopantdistribution in a sense that they cause inhomogeneous distri-butions of both. The inhomogeneous free hole concentrationdistributions as well as stress distributions together with thepresence of extended lattice defects in the seed layer of theTFSC all have a negative effect on the overall solar cellefficiency. Therefore, having a robust method that can simul-taneously determine the distribution of the free hole concen-tration and stresses in the layer quantitatively can help infinding out recipes that create material which keep thesestresses and free hole concentration variations to a minimum.Optimizing the seed layer fabrication is vital for the entireperformance of the solar cell as lattice defects can extendfrom the seed layer to the epitaxial layers grown during fur-ther processing.

IV. CONCLUSIONS

Raman spectroscopy is used to simultaneously determinelocal stress states and free carrier concentration distributionsin thin ��500 nm� multicrystalline silicon layers on glass.

First, calibration curves correlating the free carrier concen-tration to the symmetry parameter q �which describes thepeak asymmetries induced by the Fano resonances that occurat high doping levels—�1019 cm−3� were generated forthree laser wavelengths. These calibration curves were thenused in quantifying the free hole concentration in the lasercrystallized seed layer used as part of a TFSC concept. Sec-ond, the peak shifts due to doping were isolated from thosedue to stress, thereby also allowing the quantification of thestress distribution in the layer. The laser crystallized seedlayer experiences stress gradients induced by the local grainboundary structure/population which were as high as�850 MPa�12 MPa. XTEM investigations confirmed thatlocal stress gradients were most likely due to lattice defectssince the laser crystallization experiment was carried out insuch a way that the overall extrinsic thermal mismatch in-duced stress should be very low, of the order of about 3 MPaat most. The free hole concentration within the seed layerwas not uniform either; it varied from 1�1018 to 1.3�1019 cm−3. Electrical measurements showing the two-point I-V curves of the seed layer demonstrated nonuniformelectrical conductivity thus confirming the inhomogeneousdistribution of doping and therefore the free hole concentra-tion within the layer. All the aforementioned effects, defects,grain boundaries as well as the variation in stress and doping,may have an adverse effect on the progressive layers that aregrown epitaxially on the seed layer in all silicon on glassTFSC technologies resulting in their performance degrada-tion.

ACKNOWLEDGMENTS

This work was supported by the Max Planck Society inthe framework of the ‘Nanostress’ project and by the Euro-pean commission in the Energy priority within the FP7project ‘High-EF’. The authors would like to acknowledgeMs. Hopfe at MPI Halle for TEM sample preparation andMr. Bochmann at IPHT for fruitful discussions.

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