Gas phase hydrogen absorption and electrochemical performance of La2(Ni,Co,Mg,M)10 based alloys

i n t e r n a t i o n a l j o u r n a l o f h y d r o g e n e n e r g y 3 9 ( 2 0 1 4 ) 2 4 2 3 e2 4 2 9

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Gas phase hydrogen absorption and electrochemical performance of La2(Ni,Co,Mg,M)10 based alloys H. Drulis a,*, A. Hackemer a, P. Głuchowski a, K. Giza b, L. Adamczyk b, H. Bala b a b

Institute of Low Temperatures and Structure Research PAS, Wroclaw, Poland Department of Chemistry, Czestochowa University of Technology, Czestochowa, Poland

article info

abstract

Article history:

The effect of M ¼ In or Al on the hydrogenation behavior of the La2(Ni,Co,Mg,M)10 alloys at

Received 15 July 2013

room temperature is presented. The ceramic like samples have been prepared by powder

Received in revised form

metallurgy route using pure Mg- and the La2Ni9xMx alloy powder precursors. XRD analysis

14 November 2013

revealed predominantly the CaCu5-type structure for all final alloys. Partial substitution of

Accepted 22 November 2013

Co by In in La2Ni8MgCo causes a slight decrease of hydrogen concentration whereas Al

Available online 19 December 2013

addition increases this parameter. The highest hydrogen concentration of 1.87 wt.% has been reached for La2(Ni8Co0.8Al0.2)Mg composition at hydrogen pressure of 10 bar. Indium

Keywords:

addition dramatically decreases the middle-plateau hydrogen equilibrium pressure from

Intermetallic hydrides

peq ¼ 0.37 bar (In-free alloy) to peq ¼ 0.06 bar (1.7 at.% In). The electrochemical performance

Pressureecomposition isotherms

of the studied materials has been characterized using chronoamperometric and chro-

Electrochemical charge/discharge

nopotentiometric techniques. The galvanostatic hydrogenation experiments at 185 mA/g

Hydrogen capacity

discharge rate revealed the largest discharge current capacity of 355 mAh/g for La2(Ni8Co0.8Al0.2)Mg alloy. The relative diffusivity factor of hydrogen ðDH =a2 Þ varies for the tested materials in the range of (2.0e5.4)$105 s1. Copyright ª 2013, Hydrogen Energy Publications, LLC. Published by Elsevier Ltd. All rights reserved.

1.

Introduction

Rare earthenickel (REeNi) based alloys are widely applied for hydrogen storage, including the rechargeable metal hydride (Ni/MH) batteries. Materials for the above applications should reveal a high hydrogen capacity, moderate hydride stability and reasonably high hydrogen absorption/desorption rates. Most of these characteristics are usually derived from the hydrogen pressureeconcentration (pec), isotherms [1,2] and electrochemical charge/discharge measurements [3,4].

The most spread and commercialized metal hydride electrodes are mainly based on AB5-type alloys. Their hydrogen capacities usually reach 300e330 mAh/g. Many methods such as optimization of composition or doping were applied to improve both the REeNi alloys discharge capacity and cycle life [5]. In our recent papers we discussed the hydrogenation properties of LaNi5xInx [6] and LaNi4.3(Co, Al)0.7xInx compositions [7]. Partial substitution of Ni by indium (x < 0.3) significantly modifies the hydrogenation behavior. Particularly, indium decreases the hydrogen equilibrium pressure making In-doped alloys very interesting materials as negative

* Corresponding author. Tel.: þ48 71 343 5021; fax: þ48 71 344 1029. E-mail address: [email protected] (H. Drulis). 0360-3199/$ e see front matter Copyright ª 2013, Hydrogen Energy Publications, LLC. Published by Elsevier Ltd. All rights reserved. http://dx.doi.org/10.1016/j.ijhydene.2013.11.092

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MH electrodes in the Ni/MH batteries. The effect of indium is especially distinct when part of nickel (3.3e6.7 at.%) is substituted by cobalt [7]. Recent investigations of Kadir et al. [8,9] and De Negri et al. [10] have shown that also the magnesium containing (RE,Mg)e Ni alloys with the general formula of RE3xMgxNi9 may serve as promising materials for hydride electrodes owing to their high hydrogen storage capacity and good electrochemical properties. Electrochemical discharge capacities of the Mg containing alloys with PuNi3-type structure are greater than those with CaCu5-type structure. For example, the capacity of 410 mAh/g was observed by Kohno et al. [11] in the system with the composition of La5Mg2Ni20Co3 whereas the capacity of 380 mAh/g was observed by Tang et al. [12] in Mg modified alloys with CaCu5-type structure. Generally, because of low magnesium boiling point (1105  C), the final composition and hydrogen storage properties of the Mg containing alloys are strongly affected by the metallurgical process used in alloys manufacturing [13,14]. The main goal of this work focuses on the relationship between the composition and both gasphase- and electrochemical charge/discharge hydrogenation for LaeNieMg type alloys obtained by the so-called sintering metallurgy. Our interest is to find Lae(Ni, Co)Mg-based alloys with a discharge capacity greater than that of LaNi5 intermetallic compound and with hydrogen equilibrium pressure lower than 1 bar.

2.

processes had been performed three e four times, the pressureeconcentration (pec) dependencies of hydrogen desorption were measured under hydrogen pressures from 20 to 0.02 bar at T ¼ 296 K. The electrochemical charge/discharge tests were carried out in a conventional three electrode cell, consisting of a powder-composite metal hydride working electrode, a reference saturated calomel electrode (SCE) and a Pt wire counter electrode, using a CHI 1140 A (Austin, Texas) workstation. The powder composite electrodes were prepared by pressing a homogenized mixture of 85 wt.% of corresponding alloy powder, 10 wt.% of PVDF and 5 wt.% of C-45 carbon black into pellets, 0.4e0.5 mm thick. The electrolyte was Ar-saturated, 6M KOH solution at a temperature of 23  C. The chronopotentiometric method was applied to determine the current capacity and exchange current density variations as a function of cycling. The electrodes were charged at a cathodic current density of 185 mA/g for 3 h and discharged at þ185 mA/g up to anodic potential of 0.6 V (vs SCE). The relative hydrogen diffusivity factors ðDH =a2 Þ, where a denotes the mean particle radius and discharge capacities as function of cycle number were determined by a multi-potential step, chronoamperometric technique. In this method electrodes were charged at Ech ¼ 1.2 V (vs SCE) for 104 s and discharged at Edisch ¼ 0.6 V (SCE) for 104 s. More details concerning experimental procedure can be found in our previously published papers [6,7,19e21].

Experimental

Five alloys of the composition of La2Ni9Mg, La2Ni8CoMg, La2Ni7Co2Mg La2Ni8(Co0.8In0.2)Mg and La2Ni8(Co0.8Al0.2)Mg were prepared by powder metallurgy using the mechanical alloying (MA) route followed by annealing. The La2Ni9xMx (M ¼ Co and Al or In) alloy precursors and Mg (99.8 wt.%) powder have been used. The precursor alloys were arc melted from the individual metals: La (99.8%), Ni and M (99.9% purity) in high purity argon gas atmosphere. As-cast precursor alloys were mechanically crushed, milled into the powders and mixed with 8.3 at.% of Mg powder. A small excess (ca 5 wt.%) of Mg was introduced into starting powder mixtures to cover partial magnesium evaporation. Then, the powder mixtures were ball milled in a Fritsch mill under argon for 5 h with the speed rate of 500 rpm and the revolution direction was being changed every 30 min. After milling, the obtained amorphous material was pressed into the pellets and sintered in 106 bar vacuum. The sintering was carried out at 800  C for 8 h, followed by a second step at 600  C for 8 h, analogously as in paper [13]. The obtained sintered ceramics were characterized by means of X-ray diffraction (XRD) using a CuKa radiation. The gas-phase hydrogen absorptionedesorption properties of the alloys were studied by the use of Sievert’s type equipment. The samples were activated in vacuum at 250  C for 1 h, cooled to 23  C and then charged with high purity hydrogen gas (99.999% H2) at pH2 ¼ 20 bar. Several complete hydrogen absorption-desorption cycles were performed prior to the peceT measurements to ensure high rate of hydrogen exchange. The cyclic examination included the hydrogen absorption at 20 bar for 1 h and then the fast hydrogen desorption with a rotary pump for next 1 h. Once these

3.

Results and discussion

3.1.

Structure characterization

X-ray diffraction (XRD) with a CuKa radiation was used to identify the phase structure and composition of the alloys. The XRD data were collected using diffractometer X’Pert PRO PANalytical. Fig. 1 shows the XRD patterns evolution obtained

Fig. 1 e Evolution of XRD patterns of La2(Ni,Co)9Mg alloy during consecutive steps of its synthesis: (a) arc-melted La2(Ni,Co)9 precursor, (b) precursor and Mg powder mixture after ball-milling and (c) final La2(Ni,Co)9Mg powder after high temperature annealing.

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for La2Ni7Co2Mg composition, as an example, at three consecutive stages used during the alloy fabrication: (a) for the arc-melted precursor, (b) after milling of Mg with precursor and (c) for the final alloy after the annealing procedure. The initial arc-melted precursor, La2Ni7Co2 exhibits (Fig. 1a), rather complex diffraction pattern. The precursor and Mg milled together in the mechanical alloying (MA) route show the pattern typical to almost amorphous material (Fig. 1b). Fig. 1c presents the X-ray pattern of a full recrystallized material after its high temperature annealing. All final experimental XRD data were analyzed with Rietveld method using the XPert HighScore Plus software. Fig. 2 gives (as an example) the XRD profiles of Rietveld analysis for one of the studied alloys with La2Ni9Mg formula. X-ray spectra for precursors were not analyzed. The phasestructural analysis proves that all studied LaeNieMg and LaeNieCoeMg samples consist of mainly (>95%) with the phases that crystallize in the CaCu5 type structure. Apart from the main CaCu5-type structure pattern there are visible tiny patterns belonging to the impurity phases of approximate La2Ni7 composition whose abundance is on the level of ca 5 wt.%. The unit cell parameters of the main component determined from XRD data analysis are summarized in Table 1. The results of structural analysis of the tested materials indicating their CaCu5-type structure are rather surprising. As it was mentioned in Section 2, following the synthesis method proposed in [13] we expected to obtain the material predominantly with PuNi3 e type structure. Besides, it is believed that LaNi5 compound does not dissolve magnesium. The Rietveld refinement of the XRD pattern of the main component in Fig. 2 showed that Mg forms solid solution in our AB5-type material by occupying some quantity of 1a positions of P6/mmm space group: La1.98Mg0.02Ni10. We cannot expect anything much different in other samples (see X-ray results in Table 1). Therefore, we restrict such Rietveld analysis for one sample only. Similar results (La0.97Ni5Mg0,03) for

Table 1 e Unit cell parameters of the main component of the tested La2(Ni,Co,Mg)10LxMx alloys from XRD analysis. Sample

Space group

Unit cell parameters, nm

La2Ni9Mg La2Ni8CoMg La2Ni8 (Co0.8In0.2)Mg La2Ni8 (Co0.8Al0.2)Mg La2Ni7Co2Mg LaNi5 ref.

P6/mmm P6/mmm P6/mmm

0.5022039 0.3981173 0.5027354 0.4000330 0.5041739 0.3998451

>95% >95% >95%

P6/mmm

0.5035055 0.3992600

>95%

P6/mmm P6/mmm

0.5033665 0.3984836 0.50125(3) 0.39873(2)

>95% 100%

a

Abundance

c

AB5-Mg doped alloys have also been reported by Li et al. [15]. To confirm the presence of the rest of Mg in the alloys the EDX analysis has been carried out. In Fig. 3 the EDX pattern registered for the sample with nominal composition of La2Ni7Co2Mg is presented. Evidently, the peak of Mg is visible at 125 keV. Average chemical composition calculated from the normalized EDX peak intensities corresponds to La17.6 Ni58.0Co13.8Mg10.6 formula (where the numbers are atomic percentages) and it is close to the target composition of La16.7Ni58.3Co16.6Mg8.3 formula unit (LaNi5-type). It is worth mentioning that similar LaeNieMg composite materials based on LaNi5þx structure have already been manufactured and their electrochemical properties studied by Tang et al. [12]. Their EDX results indicated that the Mg content in the regions of LaNi5eMg- contained solid solution reaches value as high as 12 at.% close to the value of 10.6 at.% estimated in our alloys. The presence of LaNi5eMg doped phase (La0.78Nd0.18Mg0.03Ni3.99Mn0.19Co0.36Al0.33) have also been reported by Ozaki et al. [16] for La0.8Mg0.2Ni3.4xCo0.3(MnAl)x composition with x ¼ 0.4.

Fig. 2 e Rietveld refinement of the XRD pattern of the La2Ni9Mg alloy. Lower plot is a difference profile between experimental and calculated lines. Vertical bars correspond to the Bragg peak positions for the constituent phases. Bars for La2Ni7 impurity phase are omitted.

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Fig. 3 e EDX pattern of La2Ni7Co2Mg alloy.

3.2. Hydrogen pressureehydrogen concentration (pec) isotherms Fig. 4 shows the pec isotherms (296 K) of hydrogen gas desorption for four representative intermetallic phases of magnesium modified materials with final La2Ni9Mg, La2Ni8CoMg, La2Ni8(Co0.8In0.2)Mg and La2Ni8(Co0.8Al0.2)Mg stoichiometry. From the crystallographic point of view, all samples had practically the same phase composition so one could expect that they possess similar hydrogenation properties. On the other hand, the precursors applied had different chemical composition because of partial substitution of Ni by Co and Al/ In additions. To follow through the role of Mg element on hydrogen absorption properties of the individual LaNi5etype alloys the hydrogenation and pec characteristics of the precursor hydride phases are additionally shown. Under hydrogen pressure of 10 bar the largest hydrogen concentration equal to 1.87 wt.% has been obtained for La2Ni8(Co0.8Al0.2)Mg composite. After the first cycle, the hydrogen capacity used to drop up to ca 10% of its initial value, depending on the alloy composition. Such a behavior indicates that part of the composite material is hydrogenated irreversibly. Thus, part of the material cannot be involved in hydrogen desorption process as long as absorption/desorption cycles are carried out at room temperature. This is likely Mg2NiH4 hydride, which appears when composite material decomposes upon hydrogenation. Generally, it is worth noting the dramatic difference in hydrogenation ability between the precursor alloys (without Mg) and magnesium modified materials. For example, the hydride capacity of La2Ni8(Co0.8In0.2) Mg final alloy is over 3 times greater than that of its La2Ni8(Co0.8In0.2) precursor. The reason for these different properties one can explain by a very complex multiphase state of the precursor alloys. The corresponding situation is illustrated in Fig. 5 by X-ray diffraction spectra for Mg-free precursor (pattern “a”) and magnesium containing alloy (pattern “b”). From the application point of view, the so-called reversible hydrogen capacity (RHC) is of great importance. In this work the RHC is assumed to be the amount of hydrogen gas (expressed in mAh/g) that can be derived from fully hydrogenated material during hydrogen isothermal desorption between two points shown in Fig. 4: one marked by CH,1 bar and

second with hydrogen pressure equal to pH2 ¼ 0. Such defined RHC parameter has a very practical meaning because from NiMH battery stability point of view the negative electrode material should exhibit maximum hydrogen equilibrium pressures no higher than 1 bar at room temperature. As it results from Fig. 4(aed) all of the tested alloys containing 8.3 at.% Mg satisfactorily fulfill this criterion. In Table 2 the mid-plateau H2 equilibrium pressures, maximum hydrogen concentration in the tested materials (arbitrary assumed to correspond to p ¼ 10 bar), hydrogen content in the alloys at 1 bar and the calculated RHC values determined from pec curves in Fig. 4 are collected. As it has been already mentioned, the RHC values presented in Table 2 were estimated from the width of plateau part of pec isotherms between 1 bar and the vacuum. Therefore, the values of RHC given in Table 2 are certainly somehow overestimated. In Table 2, the Qel,disch values determined electrochemically (see Section 3.3) are also collected for comparative purpose. In practice, the values of RHC parameter reflect the expected discharge current capacities of corresponding hydride electrode for an “open” Ni/MH battery.

3.3.

Electrochemical hydrogenation

In Fig. 6 the anodic current densities (in logarithmic scale) versus discharge time recorded for cathodically charged La2Ni8(Co0.8In0.2)Mg electrode are presented for 7 successive cycles, as an example. Similar dependencies were obtained for La2Ni8CoMg, La2Ni8(Co0.8Al0.2)Mg and La2Ni7Co2Mg electrodes. Integration of anodic current density of hydrogen oxidation over the entire range of discharge time allows to evaluate the changes of discharge capacity (Qdisch) at subsequent cycles. The calculated discharge capacities vs cycle number for all of the tested electrodes are shown in Fig. 7. As it can be seen, the tested alloys usually need 2e4 cycles to reach their maximum capacity. The lowest current capacity exhibits the La2Ni8CoMg alloy (8.3 at.% Co) e its maximum value of 280 mAh/g corresponds to fourth cycle. Partial substitution of Co by In or Al (1.7 at.%) in this material is prone to distinct increase of hydrogen absorption ability. For La2Ni8(Co0.8In0.2)Mg electrode the discharge capacity was 340 mAh/g (3e4 cycle) whereas for La2Ni8(Co0.8Al0.2)Mg the capacity was as large as 367 mAh/g (2e3 cycle). Similarly great capacity (344 mAh/g for 4e5 cycle) was observed for Co-rich alloy (16.7 at.% Co) of La2Ni7Co2Mg composition. From Fig. 7 it appears that capacities of cobalt-, cobalt/indium- or cobalt/aluminum substituted La2Ni9Mgbased alloys are generally comparable to each other (ca 336 e 355 mAh/g). The electrochemical discharge capacities of LaNi5 e based alloys without Mg component used in commercial Ni/ MH batteries are usually on the level of 300 mAh/g. Thus, some of LaNi5 -Mg -based alloys reported in this paper can be considered as potential candidates for negative electrode materials in the rechargeable NieMH batteries. From the slope of the linear segments in Fig. 6 it is possible to estimate the effective coefficient of the hydrogen diffusion DH in the electrode using the following equation, which is valid for sufficiently long discharge times [17]:   6FD p2 DH ðC0  CS Þ  t logi ¼ log  2 da 2:303a2

(1)

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Fig. 5 e XRD patterns for La2Ni8(Co0.8In0.2) precursor and the final La2Ni8(Co0.8In0.2)Mg alloy.

where i denotes the measured anodic current density, DH e the effective hydrogen diffusion coefficient, d e density of the alloy, a e average radius of the alloy particles, Co e the initial hydrogen concentration in the alloy, Cs e the surface hydrogen concentration and t the actual discharge time. The sign  in Eq. (1) corresponds to the charge () or discharge (þ) processes. Because, it is hard to determine the real average particle size with satisfactory accuracy (and, thus its mean diameter a) we use the DH =a2 fraction to evaluate hydrogen diffusivity within the electrode material. The calculated values of DH =a2 (we name them “relative diffusivity factors”) are presented in Fig. 8 for subsequent cycles. As it is shown in Fig. 8 the DH =a2 values determined from the chronoamperometric measurements are on the order of 105 s1. The chronopotentiometric method is also very useful for determination of exchange current density of H2O/H2 system for hydrogen storage material as a function of cycle number [18]. According to [20,21] the exchange current density ioH2 O=H2 can be obtained from the following relationship: 1 DE logioH2 O=H2 ¼ logðia jic jÞ  2 2b

(2)

where ic, and ia, are the charge- and discharge current density, DE is the potential jump that occurs during external current switching from negative into positive values, and b e Tafel slope of cathodic/anodic straight line for hydrogen electrode (equal to 0.12 V at room temperature). As seen from Fig. 9, the H2O/H2 exchange current density increases with cycle number with certain tendency to settle down after 7e8 cycle. Only for La2Ni8Co0.8Al0.2Mg electrode we can see a progressive increase of the exchange current density with cycling. An increase in the exchange current density with cycling reflects the charge transfer rate increase at the interface between MH alloy powder and the electrolyte. The most

Fig. 4 e Hydrogen desorption isotherms of (a) La2Ni9Mg (b) La2Ni8CoMg, (c) La2Ni8Co0.8In0.2Mg and (d) La2Ni8Co0.8Al0.2Mg hydrides and their precursor hydride phases at T [ 296 K.

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Table 2 e Hydrogenation parameters of the tested La2Ni9Mg material partly substituted with Co and In/Al for nickel, determined from the pec measurements (23  C). Alloy composition

a

peq [bar]

La2Ni9Mg I/86s La2Ni8CoMg I/98s La2Ni7Co2Mg I/109s La2Ni8Co0.8In0.2Mg I/101s La2Ni8Co0.8Al0.2Mg I/106s a b c d e

b

CH,10 bar [%wt]

0.91 0.43 0.60 0.45 0.63

1.63 1.85 1.75 1.71 1.87

c

CH,1 bar [%wt] 1.29 1.52 1.34 1.37 1.67

d

e

RHC [mAh/g]

347 409 360 369 449

Qel,disch [mAh/g]

Galvan.

Chronoam.

314 230 325 339 355

280 344 340 367

Equilibrium pressure of H2 measured in the middle of plateau of pH2 ¼ f(cH) isotherm. Hydrogen concentration absorbed by the fresh sample (first cycle at pH2 ¼ 10 bar). Average hydrogen concentration in a sample when hydrogen gas pressure is equal 1 bar. Reversible capacity read from pec isotherm along plateau between pH2 ¼ 1 bar and hypothetical vacuum. Discharge capacity from galvanostatic- (at 0.5C/þ0.5C rate) and chronoamperometric measurements.

cycle number

i, mA/g

5, 6, 7

100

4 3 2

10

1

0

2000

4000

Ni Ni Ni Ni

CoMg Co In Mg Co Al Mg Co Mg

6

4

2

La 2Ni 8Co 0.8 In 0.2 Mg

1

La La La La

8

(D/a2) x 10 5, s -1

1 2 3 4 5 6 7

1000

6000

8000

10000

t, s Fig. 6 e Chronoamperometric curves of La2Ni8Co0.8In0.2Mg electrode at L0.6 V (SCE) for 7 subsequent cycles.

Fig. 7 e The discharge capacity of the studied electrodes vs cycle number determined by chronoamperometric method.

4

6

8

10

cycle number Fig. 8 e Relative hydrogen diffusivity factors of the tested electrode materials versus cycle number.

Fig. 9 e Exchange current density of the H2O/H2 system for the tested electrode materials versus cycle number.

i n t e r n a t i o n a l j o u r n a l o f h y d r o g e n e n e r g y 3 9 ( 2 0 1 4 ) 2 4 2 3 e2 4 2 9

feasible reason of this increase seems to be a continuous development of active surface with cycling. The tendency for ioH2 O=H2 to stabilize may result from appearance of corrosion products (oxide phases) at individual particles that inhibit both hydrogen transport and charge transfer at the interfacial areas.

4.

Conclusion

This paper confirms the earlier observations that the final composition and hydrogen storage properties of the Mg containing REeNi based alloys are strongly affected not only by the metallurgical process used in the alloys manufacturing but also by the subtle details of their synthesis. Nevertheless, from the application point of view the so-called reversible hydrogen capacity (RHC) is of great importance. The RHCs estimated in this paper, are quite close to the Qel,disch values measured directly from electrochemical experiments. The analysis of the mentioned RHCs shows that apart from La2Ni8CoMg the best hydrogen (both gas-phase and electrochemical) desorption performance exhibit alloys with stoichiometry of La2Ni8(Co0.8Al0.2)Mg and La2Ni8(Co0.8In0.2)Mg i.e. those with part of cobalt (1.7 at.%) substituted by Al or In. Exchange current densities of H2O/H2 system increase with cycling. The greatest exchange currents (>70 mA/g) have been found for Al-doped alloy. Established structure- and hydrogenation properties of the described La2(Ni,Co,Mg,M)10-type composites will allow a better selection and composition optimization in further development of the NieMH battery negative electrode materials with improved electrochemical performance. This optimization includes further substitutions and examination of their synergistic effects and will be a subject of our prospective investigations.

Acknowledgments The work was supported by Wroclaw Research Centre EITþ within the project “The Application of Nanotechnology in Advanced Materials” e NanoMat (POIG.01.01.02-02-002/08) cofinanced by the European Regional Development Fund (Operational Programme Innovative Economy, 1.1.2).

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