Spark-plasma-sintered porous electrodes for efficient oxygen evolution in alkaline water electrolysis

Electrochimica Acta 317 (2019) 128e138

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Spark-plasma-sintered porous electrodes for efficient oxygen evolution in alkaline water electrolysis €cker b, Stefan Loos b, Marcus Vogt b, Thomas Rauscher a, Christian Immanuel Berna € ntzsch b, * Bernd Kieback a, b, Lars Ro a b

€t Dresden, Helmholtzstraße 7, 01069, Dresden, Germany Institute of Materials Science, Technische Universita Fraunhofer Institute for Manufacturing Technology and Advanced Materials IFAM, Branch Lab Dresden, Winterbergstraße 28, 01277, Dresden, Germany

a r t i c l e i n f o

a b s t r a c t

Article history: Received 6 February 2019 Received in revised form 30 April 2019 Accepted 20 May 2019 Available online 23 May 2019

Porous electrodes for alkaline water electrolysis were prepared by spark plasma sintering, a short-time sintering technique, in combination with a space holder method. After removal of the space holders, highly porous layers of polycrystalline Ni and of a nanocrystalline Ni-Fe alloy were obtained on a metallic substrate. Both porosity and thickness of the electrocatalytic layers can be controlled by the space holder volume content and the sintering process conditions, for example, the applied pressure and temperature. The active surface of the electrode can be increased significantly by a roughness factor of up to 1,120 determined by double layer charging. The porous layers are efficient towards oxygen evolution reaction (OER), whereas activity is greatly influenced by the chemical composition. The porous Ni-Fe electrodes exhibit an extremely low OER-overpotential of 230 mV at 0.3 A cm
2 in highly concentrated KOH (29.9 wt.-%) at 333 K. Beside the high surface area, the efficiency of the porous Ni-Fe layer is characterized by a high intrinsic activity resulting in a low Tafel slope of around 23 mV dec
1 at low and 50 mV dec
1 at high current densities as well as a high turnover frequency (TOF) of approximately 3.4 s
1 at 0.3 V. The porous Ni electrodes have a lower intrinsic activity with higher Tafel slopes and lower TOF. Moreover, an excellent stability and activity under realistic operating conditions of intermittent electrolysis (up to 1 A cm
2) for 100 h was proven for the porous Ni-Fe electrode. © 2019 Elsevier Ltd. All rights reserved.

Keywords: Oxygen evolution reaction Porous electrode Alkaline electrolysis Spark plasma sintering Space holder

1. Introduction Our global energy system is facing significant challenges to reach the ambitious goals of the reduction of CO2 emissions (Paris Agreement). As the use of renewable energy such as wind and solar continues to increase, energy storage and conversion technologies are urgently needed. Water electrolysis is one of the most promising technologies to fulfill this demand. Efficient and cost-effective hydrogen production requires advanced electrolyzers with modern cell designs [1], improved membrane or separator structures as well as highly active and stable electrode materials that do not contain any critical materials such as precious metals [2]. The efficiency of electrochemical water splitting is mainly limited by sluggish reaction kinetics of the oxygen evolution reaction (OER) resulting in a high overall cell voltage [3e5]. There are several approaches to develop active and stable catalysts for the OER [4].

* Corresponding author. € ntzsch). E-mail address: [email protected] (L. Ro https://doi.org/10.1016/j.electacta.2019.05.102 0013-4686/© 2019 Elsevier Ltd. All rights reserved.

Recent investigations were mainly devoted to fundamental questions regarding the active sites [6,7] the structure and structural changes of catalytic layers during the OER [8] as well as the effect of impurities in the electrolyte on the catalyst activity. Furthermore, it was found that especially Fe impurities improve the OER activity significantly [9e14]. Some studies aimed at identifying the catalytically active sites to improve efficiency and stability of the catalysts [15,16]. However, the electrocatalysts are often prepared as thin films on 2D substrates. The OER activity can additionally be enhanced by increasing the surface area [4]. In this regard, porous and 3D electrode structures can fulfill these requirements. Furthermore, a well-defined porosity for gas transport, a mechanically stable structure and a good electrical conductivity are essential for highly active and long-term stable catalysts. In alkaline solution non-precious, earth-abundant and inexpensive transition metal-based materials can be used as catalysts [9,15,16] which often form (oxy)hydroxides on the surface before or during the OER. In this context, Ni-based catalysts have gained interest because of the high activity and stability. The OER activity

T. Rauscher et al. / Electrochimica Acta 317 (2019) 128e138

can be enhanced by incorporation of Fe into the layered Ni (oxy) hydroxide structure. Particularly, Ni1-xFexOOH with 15% to 50% Fe are among the most active catalysts towards the OER [11,14,17,18]. Furthermore, Fe impurities (1e10 ppm) that are present in reagentgrade KOH solution can be incorporated into the (oxy)hydroxide [12] forming an active Ni1-xFexOOH layer [10,14,19]. Numerous studies deal with manufacturing technologies to produce Ni-based electrodes with a large surface area or a highly porous structure, such as laser processing technologies [20,21], galvanic deposition [22] or the well-known production of Raney-Ni [23,24] by selective leaching of specific phases. Furthermore, it is known that sintering technologies can be used to fabricate porous coatings. For example, in 1981 Hall et al. [25] investigated porous sintered Ni and Ni-Fe electrodes with an enhanced surface area showing that these coatings are suitable electrodes for the OER. Rausch and Wendt [26] produced sintered Ni electrodes and compare their activity towards the hydrogen evolution reaction with Raney-Ni. Balej investigated mixed oxide electrodes in comparison to Raney-Ni anodes [23]. The oxide electrodes were fabricated by different technologies, amongst others by pressing and sintering. Beside Ni, nanocrystalline Ni-Fe-based materials exhibit good activity towards the OER [27]. In the 1980's Kreysa et al. [28] studied different amorphous/nanocrystalline (nc) alloys towards the OER and found that in addition to Co50Ni50Si15B10 different NiFe-based alloys (e.g. Fe39Ni39Mo2Si12B8) show excellent activity for the OER. Here we present a novel approach to produce porous, mechanically stable catalyst layers on a Ni substrate using spark plasma sintering (SPS), a short-time sintering technology [29,30], in combination with a space holder method [31]. As starting material, a nanocrystalline (nc) Ni39Fe39Mo4Si12B6 powder and common Ni powder were employed. This nc Fe-Ni alloy shows a good OER activity under realistic operating conditions, as shown previously [27]. Space holder techniques are well-known to fabricate porous materials for different applications [31e36]. In particular, shorttime sintering technologies such as SPS (also called field-assisted sintering) are suitable to produce advanced materials like nanostructured magnets [37], materials of biomedical materials [33,38] or lightweight components [36]. By adjusting the parameters for the sintering process (T, p, space holder volume etc.) porous layers with tailored porosity and pore size can be realized. Moreover, the SPS manufacturing process is fast, scalable, resource-saving and environmentally friendly. This study shows for the first time that long-term stable porous Ni-Fe layers with a high OER activity can be manufactured using SPS in combination with the space holder technique. 2. Experimental

129

more detail elsewhere [27]. The milled powder of nc Ni39Fe39Mo4Si12B6 has a particle size (number weighted) between 1.5 mm and 91.5 mm (D50 ¼ 15.5 mm) and was sieved below 40 mm. This sieving step was introduced to realize a homogenous mixture between the metallic powder particles and the space holding particles. The particle morphology of the different powders is depicted in Fig. 1. Ni powder has a spherical shape. In contrast, nc Ni39Fe39Mo4B12Si6 powder and the K2SO4 powder have an irregular shape. A mixture of 30 vol.-% of K2SO4 and the metallic powder (Ni or nc Ni39Fe39Mo4B12Si6) was mechanically blended using a tubular mixer (72 rpm for at least 1e5 min) to ensure a homogeneous distribution of the particles. This step was necessary to realize stable interconnections between the metallic particles as well as the Ni substrate after the sintering process. The mixture was consolidated and connected with a Ni sheet as substrate (thickness around 1 mm, 99.98% purity) using a SPS/FAST device from FCT Systeme GmbH. The punches and the die were made of graphite. Additionally, graphite foil was placed horizontally and vertically between the sample and the graphite tool. SPS processing was conducted under vacuum (~5∙10
2 mbar) with an applied uniaxial force of 4e15 kN (sample diameter: 20 mm, pressure of 12.7e47.8 MPa). Immediately after reaching the maximal temperature (measured with a pyrometer) the sample was cooled down under vacuum. More details can be found in the supporting information (Fig. S 2). The sintering conditions are summarized in Table 1. After the sintering process the space holder (K2SO4; solubility in water of 120 g L
1 at 25  C [41]) was removed by placing the SPSprocessed sample in deionized water at temperatures between room temperature and 37  C. The solvent was periodically refreshed, and the dissolution process was evaluated as described in more detail the supplementary material. For all prepared porous layers about 85e100% of the space holder was removed after 8 h of immersion suggesting that less than 15% of K2SO4 remained in the porous structure (Fig. S 3.). This indicates that some isolated K2SO4 particles would be trapped in the porous layer. The diffusion of fresh solvent and the transport of the dissolution products though the porous structure affects the dissolution time, as discussed in more detail by Arifvianto et al. [33]. Note, that no dissolution of the transition metals or metalloids was observed. After removing the space holders, the porous SPS-processed sample was cut in pieces for electrochemical and metallographic investigations. For the electrochemical studies squared specimens (~5.5  5.5 mm2) were prepared. In case of the porous nc Ni39Fe39Mo4B12Si6 layer an unwanted oxidation was prevented by using a water-free cutting oil. Electric contact was ensured from the back side of the Ni substrate using an Ag-coated Cu wire and a conducting silver epoxy resin (E-Solder, Epoxy Products). Afterwards, the connection point and the Ni substrate were covered with a KOH stable coating layer (Wepelan Plating Resist, Peters).

2.1. Materials 2.2. Structural investigations The porous electrodes were prepared using SPS (also called field-assisted sintering technique, FAST) in combination with a space holder technique. Homogeneous mixtures of space holding particles and Ni or Ni39Fe39Mo4Si12B6 powder were fabricated and sintered onto a substrate (cf. Fig. S 1). Sieved K2SO4 powder was selected as space holding particles (99% purity, particle size: 40e100 mm, purchased from VWR). Pure Ni powder with a particle size (number weighted) between 2.7 mm and 16.0 mm (D50 ¼ 6.9 mm) was chosen. Furthermore, nanocrystalline (nc) Ni39Fe39Mo4Si12B6 with a homogeneous element distribution was fabricated by melt-spinning as described in our previous works [39,40]. The melt-spun ribbons were chopped and subsequently milled (swing-milling and ball-milling procedure), as described in

The geometric density of the porous layer was determined to calculate the mean porosity of the porous layer. For this purpose the mass of the powders, the mass and the dimension of the SPSprocessed sample as well as the theoretical density of 8.908 g cm
3 for Ni and the experimentally determined density of 8.13 g cm
3 for Ni39Fe39Mo4B12Si6 were considered (Table S1). The SPS-processed samples were prepared for metallographic cross-section analysis by a grinding and polishing procedure. The samples were analyzed by light microscopy (Reichert MEF4A) and scanning electron microscopy (SEM) including energy-dispersive X-ray spectroscopy (EDX) from Zeiss (DSM 950, Noran-system 7EDX with SDD-detector). Furthermore, the microstructure of the

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Fig. 1. SEM micrographs for (a) Ni powder (b) nc Ni39Fe39Mo4B12Si6 powder and (c) K2SO4 powder used as space holder.

Table 1 Sintering parameter, porosity and thickness of porous SPS-processed layers in combination with the space holder technique (sample diameter ¼ 20 mm). sample

powder

TSPS,max/  C

pIST/ MPa

global porosity/ %

layer thickness/ mm

SPS-Ni1 SPS-Ni2 SPS-NiFe1 SPS-NiFe2

Ni Ni Ni39Fe39Mo4B12Si6 Ni39Fe39Mo4B12Si6

900 800 650 650

12.7 12.7 22.3 47.8

39(7) 51(4) 55(2) 53(4)

553(44) 782(85) 998(72) 874(21)

porous layer was characterized by X-ray diffraction Bruker D8 Advanced with a LynxEye iris detector and Cu-KaI radiation. The diffractometer delivers good signal-to-noise ratios even for Fe containing samples despite X-ray fluorescence. The diffraction analysis software “Diffrac.eva, V3.1” was used for the phase analysis [42]. More details concerning the in-house developed software can be found in [43]. The Rietveld refinement was performed using TOPAS 4.2. 2.3. Electrochemical investigations Electrochemical experiments were carried out in 29.9 wt.-% KOH at 333 K using a three-electrode arrangement where the working electrode was held in a hanging meniscus configuration. A wrapped platinum wire (~8 cm2) was used as counter electrode. A saturated Ag/AgCl electrode (AgjAgCljsat'd KCl) separated through a Haber-Luggin-capillary was utilized as reference electrode. All potentials are referred to this reference electrode (E ¼ 0.197 V vs. NHE at 25  C). The temperature of an electrochemical doublewalled cell (~400 ml) was controlled by an external thermostat (DT ¼ ±2 K). The electrolyte solution was purged with N2 before (at least 30 min) and during the experiments. Unless otherwise mentioned, all current densities are referred to the geometric surface area of the electrodes (jgeo), which was determined from a topview micrograph. The surface area was electrochemically determined by the double layer capacitance, Cdl, using cyclic voltammetry (CV) and electrochemical impedance spectroscopy (EIS). CV was performed in a potential range of ±0.05 V around the open circuit potential (OCP) with a scan rate from 1 V s
1 to 0.02 V s
1 in decreasing direction according to previous works [15,20]. In case of the Ni-Fe electrodes a potential range was chosen where no significant Faradaic processes occur (see supporting material, Fig. S 5c). The subsequent potentiostatic EIS measurement was conducted with a delay of 60 s in a frequency range between 100 kHz and 0.1 Hz or 0.01 Hz with an AC amplitude of 5 mV. The complex nonlinear least-squares (CNLS) fitting routine of the software “Gamry Echem Analyst” (Gamry Instruments) was used for data analysis. Before the CV and EIS measurement, the sintered electrodes were pretreated galvanostatically with a current density of
0.5 mA cm
2 for 10 min in order to remove surface oxides which affect the Cdl  et al. [44]. Subsequently, cyclic estimation, as mentioned by Grden voltammograms were performed to remove any adsorbed

hydrogen. Thus, the following sequence was performed: 1) GS (-0.5 mA cm
2 for 10 min); 2) CV pretreatment (
1 to
0.4 V or 0 V); 3) OCP or potentistatic measurement, CV area (±0.05 V around OCP or E ¼
0.45 V), EIS area (±0.05 V around OCP or E ¼
0.45 V with 100 kHz to 0.1 Hz or 0.01 Hz). Steady-state polarization curves were galvanostatically recorded in the range of jgeo ¼ 0.316 A cm
2 to 0.1 10
3 A cm
2 (each step for 20 s in the direction of decreasing current density). A steadystate is confirmed by a subsequent potentiostatic measurements (cf. Fig. S 8). Previously, the sample was polarized at a current density of 0.3 A cm
2 for 1 h. The overpotential, hOER, was determined by hOER ¼ E - Eeq where E is the potential measured versus the reference electrode and Eeq is the equilibrium potential of the OER which was calculated based on the Nernst equation (Eeq ¼ 0.163 V in 29.9 wt.- % KOH at 333 K). All electrochemical experiments were carried out with a Gamry Reference 3000 and were IR-drop corrected. The electrochemically active surface area (ECSA) was determined by integration of the area under the redox peak of the Ni(OH)2-to-NiOOH transition (or vice versa) [9,45] assuming that the redox active Ni sites are equal to electrochemically active centers. The voltammograms were recorded with 10 mV s
1 in a potential range of 0 V to 0.55 V after the O2 evolution at 0.3 A cm
2 for 1 h. The turnover frequency (TOF) was determined based on the charge, q, related to the conversion of NiOOH-to-Ni(OH)2 according to TOF ¼ j/ 4 q [46], where j is the current density at an overpotential of 0.3 V. The charge was obtained by integration of the cathodic peak because the anodic peak of the porous Ni-Fe electrodes overlaps with the OER (cf. Fig. S 11). Long-term measurements at different current densities (0.3 A cm
2 and 1 A cm
2) for around 100 h were performed in a two-compartment cell. The half cells were separated from each other by a separator (ZirfonPerl, Agfa Gevaert) and the water consumption was continuously compensated with a programmable injection pump (Landgraf Laborsysteme HLL GmbH). More details can be found in our previous work [27]. 3. Results and discussion 3.1. Structural characterization Fig. 2 depicts the cross-section micrographs of the porous layer

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131

Fig. 2. Optical micrographs of porous SPS-processed layers revealing a bimodal pore structure (large pores ¼ previous space holder, small pores ¼ incomplete densification/sinter porosity): (a) SPS-Ni1, (b) SPS-Ni2, (c) SPS-NiFe1, (d) SPS-NiFe2. SEM micrographs within the frame show the corresponding sinter porosity at higher magnification. The sintering conditions are summarized in Table 1.

produced by SPS processing in combination with the space holder after subsequent removal of the space holding particles. It is clearly visible that different porous structures with a pore hierarchy can be fabricated, i.e. structures containing large and small pores. Obviously, the large pores (size > 51 mm) are related to the removed space holder (K2SO4, size ¼ 40e100 mm). The small pores correspond to the incomplete densification process during the fast sintering (so-called “sinter porosity”). The latter can be adjusted by the SPS parameters (e.g. temperature, pressure). Consequently, porosities from 39% to 55% with the same space holder content of 30 vol.% can be realized as shown in Table 1. Comparing the measured porosities with the space holder content indicates that the sinter porosity is in the range between 9% and 15%. In general, an open porous structure (percolation) can be considered if the porosity is higher than 5e10% [47]. The shape of the large pores is determined by the shape of the space holding particle. In contrast, the pore structure of the small pores is presumably defined by the sintering conditions (pressure, temperature). In particular, a more uniform pore structure of elongated and even spherical pores can be found of the porous SPS-Ni layer sintered at 900  C. At 800  C a more irregular pore structure is obvious. This effect is more pronounced for both SPS-processed Ni-Fe layers. In this case the sintering porosity is affected by the applied pressure. If the pressure is decreased from 47.8 MPa to 22.3 MPa the sintering porosity is enhanced by 2%. Furthermore, pores with various shape and size are visible for the SPS-NiFe1 layer produced with reduced pressure (Fig. 2c). According to the sintering theory [25,48,49], the contact area between adjacent particles rises with enhancing densification, which results in a lower sinter porosity. Thus, the size of pores decreases and the shape becomes more spherical (intermediate sinter stage) [49], as clearly seen for the porous SPS-processed Ni layers (Fig. 2a and b). Apart from a transformation to more spherical and smaller pores, a higher fraction of closed pores can be expected with increasing densification process. Closed pores do not contribute to the accessible surface area. For the porous SPSprocessed Ni-Fe layers a more heterogeneous, open porous network is obvious (Fig. 2c and d). Here, the irregular morphology of the starting K2SO4 particles is mainly preserved indicating the

initial state of the densification process, whereas no closed porosity is expected [49]. Furthermore, the particle size distribution of the starting powder material affects the porous structure of the sintered layer. The greater and broader size distribution of the Ni-Fe powder (particle size: 1.5e40 mm) compared to the Ni powder (particle size: 2.7e16.0 mm) causes the more heterogeneous pore network in combination with the used sintering conditions. Accordingly, a more open porous network is obtained for the porous Ni-Fe layer. In summary, an incomplete densification process (initial or intermediate sintering state) as well as the greater and broader distribution of the staring powder particles leads to a higher sinter porosity, larger pores and a more open porous network. Thus, for porous layers with a highly accessible surface area an initial or intermediate state of sintering is required [49]. In combination with a space holder a bimodal pore structure was generated (pore hierarchy). The larger pores should be beneficial for an unhindered gas bubble release. Furthermore, the phase formation during SPS-processing of nc Fe39Ni39Mo4B12Si6 powder was investigated by XRD measurements. Fig. 3 depicts the recorded XRD patterns of the SPSprocessed Ni-Fe samples in comparison with the metallic powder before the short-time sintering process. The diffractogram of the as-prepared Fe39Ni39Mo4B12Si6 powder shows one strong broad peak at ~43.8 2q and the diffuse broad peaks with less intensity at high angles (~75.4 2q and ~80.4 2q). Additionally, minor reflections with less intensity were found (marked with asterisks). The broad peak indicates a nc structure, as described in our previous work [27]. The comparison of this diffractogram with the Xray pattern recorded after SPS processing indicates that grain growth and formation of different phases occur during the shorttime sintering. SPS-processed Ni-Fe samples fabricated with different pressures (22.3 MPa, 47.8 MPa) show similar diffraction patterns. This in turn means that the applied pressure does not affect the crystallinity of the porous layer, if the other sintering parameters (temperature, heating rate, dwell time) are kept constant. The reflections at ~43.9 , ~51.1, ~75.2 2q can be attributed to a solid solution of a Ni-Fe phase (fcc-structure, space group Fm3m)

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Fig. 3. X-ray diffraction patterns of the SPS-processed Ni-Fe layers (Ni39Fe39Mo4Si12B6) produced with different pressures (22.3 MPa, 47.8 MPa) in comparison to the asprepared powder. The X-ray powder diffractogram of the milled powder was published previously [27]. Reflections marked with a diamond belong to unknown phases formed in small amounts during SPS-processing. The given numbers stem from PDF4 database entries, e.g. 04-011-9007.

which is the majority phase. Furthermore, different minority phases can be found, an Fe-Si phase (trigonal, P3m1), metallic Mo (bcc, Im3m) and probably a Mo-Ni carbide phase (W6Fe6C-type structure, Fd3m). Some diffraction peaks cannot be assigned to any phases mentioned above (marked with a diamond). By performing a robust Rietveld refinement and disregarding the unassigned reflections, (cf. Fig. S 9) the fraction of the major phase (solid solution of fcc-Ni-Fe) was estimated to be ~84% (w/w) with a crystallite size of 9.8(2) nm. To summarize, the sample after SPS-processing contains a Ni-Fe solid solution as majority phase with a crystallite size of approximately 10 nm (cf. Fig. S 9). A homogeneous distribution of Fe can be expected indicated by EDX analysis (Fig. S 10). The formation of a solid solution of Ni and Fe indicates that an incorporation of Fe in the Ni-hydroxide layers is likely which is a necessary condition for a significant improvement of catalytic OER activity [9,14,17]. 3.2. Electrochemical characterization 3.2.1. Real surface area The real surface area of the electrodes was determined by CV and subsequent EIS measurements at the OCP according to McCrory et al. [15] and our previous work [20]. For the Ni-Fe electrodes, a potential of
0.45 V was chosen where no significant Faradaic processes were expected (see supporting information, Fig. S 5c). For all electrodes the capacitance was determined by CV and EIS measurements as shown exemplarily in Fig. 4. The impedance spectra of the porous electrodes show a depressed or stretched semi-circle and, in most cases, an additional 45 line at high frequencies. The latter is typical for porous electrodes [51,52]. Accordingly, the Bode representation reveals a minimum in the phase at low frequencies (LF) and an additional

capacitive contribution at high frequencies (HF). Therefore, two time constants (serial arrangement of RC-terms) are considered similarly to previous works [20,50,51]. Representative CVs and impedance spectra of all electrodes and more details can be found in the supporting information (cf. Fig. S 4 e S 7). The double layer capacitance and the roughness factor, Rf ¼ C/Csmooth [53,54], determined by scan rate dependent CVs and EIS are summarized in Table 2. In principle, both methods tend to agree and the deviation between the calculated roughness factors is less than 30%. For the sake of simplicity, the capacitance and the Rf values determined by CV will be discussed hereinafter. The capacitance of the smooth and the sandblasted Ni electrodes are in line with previous works [20,50,55]. Sandblasting of Ni leads to a roughening of the surface by the factor of about 3. For porous SPS-processed Ni electrodes a surface roughness factor of 69e190 is obtained depending on the porosity and the porous network of the layer. For these layers, the value of Rf rises by the factor of 3 roughly, when the porosity increases from 39% (SPS-Ni1) to 51% (SPS-Ni2). It should be noted, that this effect is related to an enhanced sinter porosity (incomplete densification process) as the content of the space holder is kept constant at 30 vol.-% (Table 1). A relatively high roughness factor was found for the porous SPSprocessed Ni-Fe layer, with 552 and 1,121, respectively. Taking into account that only a slightly higher mean porosity is determined for SPS-processed Ni-Fe porous layers (cf. Table 1), the considerable surface enhancement is related to the heterogeneous pore network and the more open porous structure of the catalytic layer as shown in Fig. 2c and d and discussed above (chapter 3.1). Especially, the lower pore size of the SPS-Ni layer gives rise to a higher fraction of closed pores that do not contribute to the electrochemically determined surface area, naturally. In short, these results demonstrate that the roughness factor crucially depends on the porosity of the sintered layer in combination with the space holder method. Especially, the nature of the porous network due to the incomplete densification process affects significantly the determined surface roughness. Accordingly, using the same space holder content the surface roughness can be adjusted from 69 to 1,121. Interestingly, the surface roughness of the porous SPS-processed Ni-Fe layers (552 and 1,121) is in the range of Rf values reported on Raney-Ni electrodes (Rf ¼ 270e12,000, cf. Table 2) and showing that highly porous coatings can be generated by using the short-time sintering in combination with the space holder method. 3.2.2. OER activity of the porous electrodes The overpotential for OER at 0.3 A cm
2 for 1 h is shown in Fig. 5a. The overpotential of smooth Ni increases initially, reaching a maximum (after 1e2 min) and decreases quickly thereafter. This is in line with our previous results [27] and is tentatively attributed to the absorption of Fe from electrolyte impurities into NiOOH as observed by several other groups [10,14,19]. The KOH used in the present work was not purified (~5.7 ppm Fe, measured with polarography) in order to simulate conditions close to an industrial application. For the porous SPS-processed Ni electrodes a slight deactivation is seen, whereas the porous SPS-processed Ni-Fe samples show a stable overpotential over 1 h. It is obvious that the porous SPS-processed Ni electrodes show a lower overpotential than sandblasted and smooth Ni. The lowest overpotential of around 230 mV at 0.3 A cm
2 was obtained for the porous SPS-processed Ni-Fe electrodes. However, no differences were observed according to the different roughness factors of 552 and 1,121 for the SPS-NiFe2 and SPS-NiFe1, respectively (cf. Table 2 and Table 3). Similar effect can be found for the porous SPSprocessed Ni electrodes. Typical steady-state polarization curves are plotted in Fig. 5b.

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Fig. 4. CV (a, b) and EIS measurements (c, d) to determine the double layer capacitance, Cdl, in 29.9 wt.-% KOH at 333 K exemplarily shown for (a, c) SPS-Ni2 and (b, d) SPS-NiFe1. For SPS-Ni2 the measurements were conducted around the OCP. In case of SPS-NiFe1 a potential of
0.45 V was selected where no significant Faradaic current were expected (cf. Fig. S 5c). CVs were performed with ±0.05 V around the applied potential with a scan rate from 1 V s
1 to 20 mV s
1. The average of the anodic and cathodic current density was plotted as a function of the scan rate to estimate the Cdl from the linear region. The impedance spectra were fitted with a modified Randles model or an equivalent circuit with two time constants according to previous works [20,50,51]. Line ¼ fit, points ¼ data.

Table 2 Double layer capacitance, Cdl, determined by CV and EIS as well as the roughness factor, Rf, of different porous SPS-processed electrodes in comparison to smooth and sandblasted Ni in 29.9 wt.-% KOH at 333 K. For the roughness factor the average capacitance of smooth Ni, Csmooth, determined by each method was considered. Literature values determined on porous Raney-Ni electrodes are shown for comparison. sample

Ni smooth Ni sandblasted SPS-Ni1 SPS- Ni2 SPS-NiFe1 SPS-NiFe2 Ni-Zn Ni-Al-Sn Ni-Al-Cr-Cu Ni66Zn14P20 Ni/Zn 50:50 wt.-% Ni-Al a b c d e

double layer capacitance, Cdl/ mF cm
2

roughness factor, Rf

CV

EIS

CV

EIS

25(4) 79(7) 1,672(142) 4,659(206) 27,357(763) 12,976(1400)

15(2) 49(5) 1,422(325) 3,465(992) 11,085(234) 6,070(352)

1.0(1) 3.2(3) 69(6) 190(8) 1,121(140) 552(109)

70,000

289,000c 120,000d 128,000e

1.0(1) 3.3(4) 70(26) 234(70) 1,283(154) 408(50) 1,200a 4,379 4,800d 6,400e 7,000e 270

5,410

3,500e

source other method

12,000(1000)b

[26] [86] [87] [88] [50] [54]

EIS was applied to gas evolving electrodes (shielding effect of gas bubbles). By anodic oxidation (galvanostatic anodic charging). At
0.1 V vs RHE. Rf was calculated based on the measured Cdl of 66 mF cmb for smooth Ni. Average value in the HER range (overpotential ¼ 0 to
0.2 V). Cdl of smooth Ni was assumed to be 25 mF cm
2. Average double layer capacitance in the HER range. Cdl of smooth Ni was assumed to be 20 mF cm
2.

For all electrodes two linear regions are observed, one at low (jgeo < 10 mA cm
2) and one at high current densities (jgeo > 31 mA cm
2) with Tafel slopes of 23e50 mV dec
1 and 50e65 mV dec
1 (see Table 3 and Fig. S 8). The lowest Tafel slope of around 25 mV dec
1 in the low current region was found for the porous SPS-processed Ni-Fe electrodes indicating a highly active electrode towards OER in this current range. For all Ni electrodes a slightly higher Tafel slope of 38e50 mV dec
1 was determined. Similar values on smooth [56] and on porous, sintered [25] Ni were reported previously under identical conditions (80  C, 30 wt.-% KOH). The increased Tafel slope of the Ni electrodes in the high current range is probably related to a change in the reaction mechanism (rate-determining step) [4] and implies a reduced intrinsic activity.

Furthermore, the Tafel curves of the porous electrodes are shifted to lower overpotentials indicating that more active sites per geometric surface area are available. In general, an enhanced OER activity with increasing surface area is likely, if the current density is referred to the geometric area instead to the real surface area [57]. From theoretical perspective, the reduction of the overpotential with enhanced surface area (Fig. 5a) can be expressed by the Tafel equation, h ¼ a þb lg j, assuming that the real current density, jreal, is defined with j ¼ jreal ¼ jgeo/R. R is the surface enhancement factor, which can be estimated by different electrochemical and physical methods [57]. Consequently, a correlation of the h value with the R value should be described by the Tafel slope, -b [20,25] (see supporting information, Fig. S 12 a). This relation will be valid if we

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Fig. 5. (a) Chronopotentiometric measurements at 0.3 A cm
2 for 1 h and (b) steadystate polarization curves (Tafel plots) for the porous SPS-processed Ni and SPSprocessed Ni-Fe samples in comparison to smooth and sandblasted Ni. The measurements were performed in 29.9 wt.-% at 333 K.

electrochemically active surface area (AECSA) describe the amount of specific catalytically active sites onto the surface. Consequently, AECSA can be significantly smaller than Areal [44], if a fraction of the surface is blocked by trapped gas bubbles (shielding effect [58]) and/or by inactive sites. In other words, the reaction takes place locally on specific catalytically active sites instead of the entire available surface area. To measure the ECSA the double layer capacitance is often used [16,57]. On the other hand, the ECSA is estimated by the charge of the voltammetric peak which correlates with the Ni(OH)2-toNiOOH transition, thereby assuming that the number of redox active Ni sites is equal to the number of electrochemically active sites for the OER. This concept is used for Ni as well as Ni-Fe catalysts [9,19,59]. To measure the ECSA based on the charge of the Ni(OH)2-to-

Table 3 Overpotential at 0.01 and 0.3 A cm
2 and Tafel parameters (exchange current density, jgeo,0 and the Tafel slope, b) for the OER in 29.9 wt.-% KOH solution at 333 K after 1 h O2 evolution at 0.3 A cm
2. The current densities are referred to the geometric surface area. sample

A

overpotential, h (0.01A/ 0.3a A cm
2)/ mV

Ni, smooth

258(1)/ 345(2)

Ni, sandblasted

259(6)/ 350(2)

SPS-Ni1

219(5)/ 304(4)

SPS-Ni2

214(6)/ 301(3)

SPS- FeNi1

168(7)/ 231(6)

SPS- FeNi2

165(2)/ 230(3)

high j region j0,


2 geo/ A cm

1.53(0.17) ∙10
06 1.55(0.12) ∙10
06 3.45(2.53) ∙10
06 5.26(3.13) ∙10
06 1.44(1.23) ∙10
05 1.21(0.83) ∙10
05

low j region b/ mV dec
1

j0,

65(1)

2.51(2.42) ∙10
08 1.10(0.67) ∙10
09 2.17(2.22) ∙10
07 3.20(1.74) ∙10
08 4.45(6.13) ∙10
09 2.36(1.89) ∙10
09

66(2) 61(3) 63(4) 53(4) 51(3)

geo/

A cm
2

b/ mV dec
1 42(8) 36(1) 46(6) 38(2) 25(5) 24(2)

Estimated from the Tafel plots. Determined by the galvanostatic measurements after 1 h.

a

assume that the entire porous layer contributes to the OER (considering no blocking effect of gas bubble) and no change of the intrinsic activity occurs. At high current densities, we observed a Tafel slope of 50e65 mV dec
1 (cf. Table 3). Therefore, a surface enhancement by one order of magnitude should lead to a reduction of the overpotential about 50e65 mV dec
1surface. Consequently, minor changes of the surface area by the factor of 2e3 leads one only to an overpotential reduction of 10e18 mV. This might be the reason for the nearly similar overpotentials and identical Tafel curves for both porous SPS-processed Ni as well as the porous SPSprocessed Ni-Fe layers. As stated above, this approach assumes that no variation of the intrinsic OER activity occurs and presupposed a valid estimation of the surface enhancement factor (cf. Fig. S 12 a, b). The intrinsic OER activity can be expressed by the turnover frequency (TOF) [13] and will be discussed below. In fact, these results show that porous catalytic layers are important to utilize large surface areas and increase the OER activity, however, the positive effect of the surface area enlargement (without affecting the intrinsic activity) on the overpotential reduction is low (50e65 mV dec
1surface). 3.2.3. Measurement of the electrochemically active surface area (ECSA) Different physical and electrochemical methods exist to determine the surface area, as discussed in detail by Trasatti and Petrii [57]. Furthermore, differences between the real and electrochemically active surface area (ECSA) are discussed [44]. The real surface area (Areal) is the total interface of the electrode, whereas the

NiOOH transition, the CVs after the oxygen evolution at 0.3 A cm
2 for 1 h are conducted as shown in Fig. 6. The CV of smooth Ni is characterized by a broad cathodic peak (C1) at 0.186 V and two well defined anodic peaks at 0.258 V (A1) and 0.313 V (A2) (cf. Table 4). The second anodic peak A2 has a lower intensity. It is reported that the anodic and the cathodic peaks are associated with the conversion from Ni(OH)2-to-NiOOH and vice versa [60e62]. According to the classical Bode scheme [63] (see Fig. S 14) two hydroxide (a/b) and oxyhydroxide (g/b) phases exist and the transitions a/g and b/b can occur. Thus, the two anodic peaks can be explained by the a/g and the b/b-transition, whereas the transition from b-Ni(OH)2 to bNiOOH occurs at higher potentials [63e65]. Thus, we assume that the first anodic peak A1 is related to the a/g and the second peak A2 corresponds to the b/b transition - in accordance with reported results on Ni [10,66,67]. It should be noted that the redox reactions of Ni are discussed controversially and some authors propose new hydroxide and oxyhydroxide phases leading to an extension of the classical Bode scheme [66,68]. The smooth Ni-Fe sample shows a rather similar CV in comparison to pure Ni, in accordance with our previous results [27] and Ni56.5Mo23.5Fe10B10 [69] electrodes; however, instead of an A2 peak a plateau is obvious. Furthermore, no significant dissolution of Mo during cycling is seen which is in line with studies on nanocrystalline Ni-Mo-B alloys with different Mo content [39]. CVs recorded on the porous SPS-processed Ni as well as SPSprocessed Ni-Fe produce a slightly different behavior related to the smooth electrodes. Compared to smooth Ni the oxidative scan of both porous SPS-processed Ni electrodes reveals that the

T. Rauscher et al. / Electrochimica Acta 317 (2019) 128e138

135

oxyhydroxide layer. Additionally, for all electrodes only one cathodic peak, C1, with a peak potential of 0.137 V to 0.192 V is detected. This is consistence with previous observations on Ni electrodes in highly concentrated KOH [66]. In summary, a mixture of b- and g-NiOOH is expected to be present after oxygen evolution, in accordance with earlier results [10]. Based on our results we conclude that with increasing surface area the amount of g-NiOOH (overcharging phase) decreases. Especially, in case of the porous SPS-processed Ni-Fe electrodes we expect that only b-NiOOH will be formed. In some publications the b-NiOOH is considered to be the “right type of oxide” for oxygen evolution [70,71] or proposed to be the active phase [66]. Yea et al. concluded that the b/b transition is more active towards OER than a/g [68]. These observations are similar to our findings if considering the overpotential at practical current densities. However, even for the pure Ni layers, the incorporation of Fe from unpurified KOH has to be taken into consideration [9,10], which leads to the formation of active Ni-Fe-layered double hydroxide structures, thereby improving predominately the intrinsic activity. In this study we used non-purified KOH to mimic quasi-industrial conditions. Therefore, we can not clarify beyond doubt that the different NiOOH phase are responsible for the observed OER activity. A discussion of the intrinsic activity will be presented below. With respect to the ECSA based on the Ni(OH)2-to-NiOOH transition (Fig. S 11), Table 4 summarizes the charge and the calculated roughness factor rq. The roughness factor is the integrated surface charge density of the NiOOH reduction to Ni(OH)2 scaled with the charge of smooth Ni [45]. The charge on smooth Ni (~3.1 mC cm
2) is in good agreement with literature data

Fig. 6. Cyclic voltammograms of (a) smooth, sandblasted Ni and smooth Ni-Fe (ribbon) and (b) the porous SPS-processed Ni and SPS-processed Ni-Fe samples. Data were recorded with a scan rate of 10 mV s
1 after the OER evolution at 0.3 A cm
2 in 29.9 wt.-% at 333 K for 1 h starting from anodic potentials to 0 V.

Table 4 Peak potential of the oxidation peak (Ep, a1; Ep, a2), cathodic peak (Ep, c1) and the charge, q c1, by integration of the peak C1. CVs are shown in Fig. 6. The roughness factor, rq, was calculated with qc1, measured/ qc1, smooth. Furthermore, the TOF was determined at an overpotential of 0.3 V based on the charge of the peak Ep, c1 according to Godwin and Lyons [46]. sample

Ep,

Ni, smooth Ni, sandblasted NiFe, ribbon SPS-Ni1 SPS-Ni2 SPS-FeNi1 SPS-FeNi2

0.186(2) 0.189(10) 0.192(2) 0.146(7) 0.137(2) 0.176(11) 0.182(2)

c1/

V

Ep,

a1/

V

0.258(4) 0.268(8) 0.272(1) 0.256(10)A 0.248(5) A not observed not observed

Ep,

a2/

V

0.312(1) 0.311(6)A plateau current 0.324(5) 0.324(1) 0.337(14) 0.327(2)

qc1 a/ mC cm
2

rq/ -

TOF / s
1

3.1(2)b 7.6(3) 5.7(18) 193(26) 508 443(17) 520(61)

1(1) 2.3(2) 1.8(6) 62(9) 163 141(4) 167(20)

4.8(5) 2.1(1) 3.8(2) 0.32(3) 0.1 3.1(4) 3.5(3)

A

Plateau or shoulder. Charge calculated by integration of the 2nd cycle (see exemplary Fig. S 11). b The integrated charge is comparable with the literature (1.5 mC cm
2 after 6 h conditioning in 1 M KOH at RT [19]; 1.3 mC cm
2 in 1 M NaOH at 25  C [59]; 2.1 mC cm
2 in pH 14 solution at RT, film thickness: 16 nm [72]). a

intensity of A1 decreases and simultaneously the peak height of A2 increases. The enhanced intensity of A2 indicates a higher amount of b-NiOOH. Considering that the CVs were recorded after O2 evolution at 0.3 A cm
2 for 1 h, this effect might be assigned to structural changes during the OER. Precisely, it is proposed that the amount of g-NiOOH (overcharging form according to the Bode scheme, Fig. S 14) is potential-dependence and at more anodic potentials a higher amount of g is likely [66,70]. The overpotential of smooth Ni after 1 h at 0.3 A cm
2 is much higher (Fig. 5a), thus, a higher amount of g-NiOOH (overcharging form) can be expected. Interestingly, for both porous Ni-Fe electrodes the oxidative scan reveals a complete suppression of A1. Only the peak A2 is visible which overlaps with the onset of the OER (comparing smooth Ni-Fe with SPS-processed Ni-Fe) due to the higher OER activity of these layers. The absence of A1 implies that the present of Fe in the bulk material influences the transformation of the

[19,45,59,72]. Slight differences might be attributed to a variation in the film thickness of the Ni oxyhydroxide layer due to the different experimental conditions. For all Ni electrodes rq (peak charge) and Rf (double layer charging) are comparable as shown in Fig. 7a. However, for the porous SPS-processed Ni-Fe layer the rq values are reduced about 3e9 times compared to the roughness factor, Rf, determined by double layer charging (CV or EIS). As a result, similar rq values were determined for both SPS-processed Ni-Fe electrodes and the SPSNi2 electrode (cf. Table 4). A lower roughness factors (rq) based on the peak charge compared to those determined by the Ni-toNi(OH)2 transition was also found for Raney-Ni electrodes by ttir et al. [45]. The authors explained the difference by Kjartansdo the poor utilization of the inner branched structure or by blocked nano-structures with hydrated oxides. Furthermore, it is necessary to take into account that only the redox active Ni sites are

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Fig. 7. (a) Comparison of the roughness factor Rf (estimated by CV and EIS) and rq (determined by integration of the redox peak) for the different electrodes. (b) Dependency of the turnover frequency (TOF, h ¼ 0.3 V) on the ECSA for the different Ni (black) and NiFe (blue) electrodes (A related to a smooth, metallic ribbon of NiFe). The black triangle highlighted in red shows the TOF of smooth Ni in purified KOH (purification procedure described in [27].). (For interpretation of the references to colour in this figure legend, the reader is referred to the Web version of this article.)

considered as active sites for the OER and, thus, Fe sites were disregarded. On the other hand, the surface area determined by the double layer charging is sometimes inaccurate because the capacitance is affected by the surface chemistry of the electrode which depends on the applied potential [57,73] and on pretreatments (etching, reduction of native surface oxides) [44]. Furthermore, the capacitance can be affected by Faradaic reactions. For instance, pseudocapacitance phenomena, caused by adsorption (e.g. upd/opd of hydrogen), intercalation or redox reaction [16,74e77], could lead to additional capacitive effects. In addition, real surfaces are not ideally polarized electrodes [57,78], which influences the determination of the double layer capacitance. In the present work, a slight discrepancy between the roughness factor based on the Ni(OH)2-to-NiOOH transition, rq, and the Rfvalue determined by the double layer charging is found. This finding might be further attributed to inactive surface regions, which could be related to the minority phases found by the XRD analysis (see XRD, 3.1 section). It is possible that these phases are inactive towards the OER. In other words, a fraction of the surface area is covered by inactive phases resulting in a lower active surface area. Thus, the oxygen evolution proceeds only on the active fraction of the surface instead of the whole accessible surface area. In this regard, the charge of the Ni(OH)2-to-NiOOH transition is used for the estimation of the ECSA (accepting the drawback that only the redox-active Ni sites are considered as active centers for the OER). 3.2.4. Intrinsic activity, turnover frequency (TOF) To compare the intrinsic activity of the catalytic layer, the turnover frequency (TOF) was calculated based on the charge related to the conversion of NiOOH-to-Ni(OH)2 and the current at an overpotential of 0.3 V according to Godwin et al. [46] (cf. Table 4). The TOF is defined as the number of O2 molecules evolved per second and active Ni site. In Fig. 7b, the TOF is plotted as a function of the ECSA for all investigated electrodes. Obviously, the TOF of the Ni electrodes decreases with increasing ECSA. Therefore, the SPS-Ni2 exhibits the lowest TOF (0.1 s
1), which is more than one order of magnitude lower than the value of smooth Ni (4.8 s
1). In contrast, the TOF (3.8e3.1 s
1) of the Ni-Fe electrodes is almost independent of the ECSA and comparable with the value of smooth

Ni. Similar or slightly higher values are reported previously for NiFe catalysts (cf. Table 5) [9,19,79e83]. However, a direct comparison is challenging due to different experimental conditions, the definition of “active sites” as well as the determination of the (electrochemical) active surface area [80,81,83]. Interestingly, the TOF for smooth Ni is similar to the reported value of 1.16(7) s
1 (at 0.33 V in 1 M KOH) for aged Ni(OH)2, [46], whereas the aging was carried out by voltammetric cycling in reagent grade KOH (Fe-containing solution). Under these conditions, the incorporation of Fe impurities into the NiOOH is likely, as shown in previous works [9,11e13] and expected by other authors [19,27]. Furthermore, reported TOFs of Ni-Fe catalysts in Fe-free KOH (cf. Table 5) reveal similar intrinsic activity compared to aged smooth Ni in Fe-containing alkaline media [10]. In consistence with this hypothesis, we measure a significantly reduced TOF of 0.1 s
1 on smooth Ni in purified KOH (Fig. 7b). This finding was further substantiated by the study of Burke et al. [13]. The authors found that 8% Fe is incorporated into the CoOOH catalyst after polarization at 350 mV for 2 h, which increases the TOF from 0.007(1) s
1 to 0.026 s
1. In Fe-free KOH, the TOF was not affected by an anodic polarization procedure [13]. Furthermore, NiOOH absorbs Fe from electrolyte impurities much faster (5% after 10 min [9,19] or 5% after 12 CVs [10]). Therefore, we attribute the significant variation of the TOF in case of the porous SPS-processed Ni electrodes to the incorporation of Fe impurities (5.7 ppm measured with polarography) into the Ni oxyhydroxide from the reagent grade KOH used in this work. Precisely, the decreased TOF with increasing ECSA is attributed to a reduction of absorbed Fe content per surface area. In other words, enhancing the surface area leads to an increased number of active Ni sites and the limited amount of Fe is spread over a much larger surface area. Accordingly, the local concentration of incorporated Fe into the NiOOH layer is much lower. Consequently, for smooth Ni (expecting a Ni1-xFexOOH layer due to the Fe incorporation) a higher TOF was determined, whereas with increasing ECSA the TOF reduces by more than one order of magnitude (Fig. 7b). The TOF of the porous Ni-Fe electrodes is nearly independent of the ECSA indicating that Ni1-xFexOOH layers with similar Fe contents can be expected. Consequently, the improved OER performance of porous SPS-processed Ni-Fe electrodes in comparison to porous SPSprocessed Ni electrodes can be ascribed to a better intrinsic activity. This finding is further substantiated by comparing the reduction

Table 5 Comparison of the TOF for different Ni based catalysts in alkaline solution with and without Fe impurities. sample

solution

TOF / s
1

TOF at h/ mV

source

Ni, smooth NiFe, ribbon SPS-Ni1 SPS-FeNi2 Ni (hydroxide) A Ni45Fe55 Au-Ni0.75Fe0.25OOH FexNi1-xO, x ¼ 10 Ni0.9Fe0.1Ox NiFe nanoparticles NiFe-LDH/CNT Exfoliated NiFe LDH

29.9% KOH 29.9% KOH 29.9% KOH 29.9% KOH 1 M NaOH 0.1 M NaOH, Fe-free 1 M KOH, Fe-free 0.5 M KOH 1 M KOH, Fe-free 1 M KOH 1 M KOH 1 M KOH

4.8(5) 3.8(2) 0.32(3) 3.5(3) 1.16(7) 2 2e0.2a 1.9 0.21(3) 6.2(1.6) 0.56 0.24b

0.3 0.3 0.3 0.3 0.33 0.3 0.3 0.3 0.3 0.3 0.3 0.3

This This This This [46] [81] [9] [82] [19] [83] [79] [80]

work work work work

Note, that discrepancies between TOFs can be attributed to the variation of the electrochemical conditions as well as the definition of the active sites (e.g. redox active sites versus total catalyst mass from in-situ QCM, ICP-OES or ex-situ XPS) [13,80,81]. The TOF based on the redox active sites is considered as upper limit [81]. A The hydrous nickel oxide film was aged by potential cycling. a TOF depends on the film thickness. b TOF was calculated based on the total charge of the Ni2þ/Ni3þ redox wave.

T. Rauscher et al. / Electrochimica Acta 317 (2019) 128e138

of the overpotential at constant geometric with surface enhancement factor based on the peak charge, as shown in the supporting information (Fig. S 12). From the practical point of view, both a high surface area as well as a high intrinsic activity (high TOF) is required to reduce the overpotential. For a high intrinsic activity, the incorporation of predefined amounts of Fe in form of electrolyte impurities into the Ni electrode should be considered. Based on our results we assume that in case of Ni the incorporated Fe content reduces with increasing ECSA, if using reagent grade KOH with a limited amount of Fe (Fe-containing solution). 3.2.5. Stability measurement Long-term stability and activity measurements at different current densities (0.3 A cm
2 and 1 A cm
2) were performed on the porous SPS-NiFe2 as it was the most active catalytic electrode in this study. Fig. 8a shows the progression of the overpotential. After 20 h at 0.3 A cm
2 an overpotential of 247 mV was measured which is slightly higher than the value after 1 h of 230 mV. During the operation at a current density of 1 A cm
2 a similar tendency is observed. However, the deactivation rate decreases with increasing electrolysis time, as indicated in Fig. 8a. After 100 h of operation at different current densities, an overpotential of 388 mV at 1 A cm
2 was measured and the overpotential approaches a nearly constant level. The overpotential of porous SPS-NiFe2 is significantly lower compared to that of smooth Ni under similar conditions (776 mV at 1 A cm
2 [27], see also stability plots in Fig. S 13). The steady-state polarization curve shows that two linear regions with a Tafel slope of 32 mV dec
1 at low and 93 mV dec
1 at high current densities can be found (Fig. 8b), which are slightly higher in comparison to the values after 1 h. In principle, an increased Tafel slope indicates a deactivation of the electrode during long-time operation. Possible reasons for the loss in activity are the decrease of the ESCA, the

Fig. 8. Stability measurement on porous SPS-NiFe2 at different current densities (step at 0.3 A cm
2 for 20 h, step at 1 A cm
2 for 5 h) for 100 h in 29.9 wt.-% KOH at 333 K. (a) overpotential versus electrolysis time, overpotentials and degradation rates at specific points are indicated. (b) Steady-state polarization curves after the stability measurements. (c) Comparison of X-ray diffraction pattern of the SPS-NiFe2 before and after the stability test and the difference curve (diff.). Yellow lines indicate the majority phase (fcc NiFe, see Fig. 3). (For interpretation of the references to colour in this figure legend, the reader is referred to the Web version of this article.)

137

decrease of the intrinsic activity and/or the reduced conductivity of the oxide film due to aging of the NiOOH film. The most probable deactivation mechanism is related to the loss of intrinsic activity and/or the reduced conductivity of the oxyhydroxide film forming during the OER. This is substantiated by the fact that no decomposition of the porous sintered layer was observed. A comparison of the X-ray patterns before and after stability measurement indicates that no changes of the surface-near region by selective dissolution of specific phases occur (Fig. 8c). Moreover, no physical damage of the sintered layer was found, in contrast to some sintered porous Ni coatings investigated earlier [25]. Consequently, the slight increase of the overpotential can be tentatively attributed to a decrease of the intrinsic activity and/or the loss in conductivity rather than to an unstable catalytic coating. This hypothesis is corroborated by previous works which propose an alternation of the microstructure [27,70,84,85]. For instance, the degradation is attributed to the conversion of Ni3þ to Ni4þ ions in the oxyhydroxide. This modification results in a reduced conductivity as well as an alteration of the active sites. More can be found elsewhere [66,70,84]. In summary, the porous sintered layer is mechanically stable to withstand the gas pressure which builds up inside the pores during O2-evolution, even at current densities up to 1 A cm
2. Accordingly, SPS processing is an appropriate technology to ensure a solid and stable connection between the individual particles of the catalytic layer. This observation suggests - in accordance with the findings above - that stable, porous and electrochemically active layers can be produced via SPS processing in combination with the space holder method. 4. Conclusions In this study, we demonstrate that highly active porous electrodes towards the oxygen evolution reaction can be produced by powder metallurgical route based spark-plasma sintering (SPS) in combination with a space holder method. The removal of the space holder offers the ability to produced stable layers with a defined porosity. The porosity of the layer can be further adjusted by the sintering conditions (e.g. temperature, force) due to the incomplete densification process which results in an additional open porous structure. Especially, the nature of this porous network defined the accessible surface area and, therefore, determined roughness factor (Rf ¼ 69e1,121, determined by the double layer charging). Furthermore, the OER activity of the porous layer is significantly affected by the composition of the layer. The highest activity towards the OER with an overpotential of 230 mV at 0.3 A cm
2 is obtained on porous, nanocrystalline Ni-Fe layer. The excellent performance is attributed to the enhanced surface area and, additionally, to a high intrinsic activity resulting in a low Tafel slope of 24 mV dec
1 at low and 51 mV dec
1 at high current densities as well as a high TOF (3.5 s
1 at 0.3 mV). The latter is attributed to the homogenous Fe-distribution in the porous layer and demonstrates that the nanocrystalline state (with crystallites in the range of 10 nm) is an excellent precursor forming an active Ni1-xFexOOH layer during the OER. The lower performance of the porous SPSprocessed Ni electrodes correlates to a lower intrinsic activity (TOF ~ 0.1 s
1). Furthermore, the long-term activity and stability of the porous SPS-processed Ni-Fe layer was proven for 100 h operation at current densities up to 1 A cm
2. These findings can certainly be transferred to other electrochemical reactions where active and porous electrode structures are needed. Acknowledgment This work was funded by German Research Foundation (DFG project code KI 516/24-1). The authors are grateful for this financial

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support. Furthermore, we thank Mrs. Lutze of the metallographic preparation of the samples. Appendix A. Supplementary data Supplementary data to this article can be found online at https://doi.org/10.1016/j.electacta.2019.05.102. References [1] R. Phillips, A. Edwards, B. Rome, D.R. Jones, C.W. Dunnill, Int. J. Hydrogen Energy 42 (2017) 23986. [2] S. Marini, P. Salvi, P. Nelli, R. Pesenti, M. Villa, M. Berrettoni, G. Zangari, Y. Kiros, Electrochim. Acta 82 (2012) 384. [3] M. Gattrell, B. MacDougall, Reaction Mechanisms of the O2 Reduction/evolution Reaction, American Cancer Society, 2010. https://onlinelibrary.wiley.com/ doi/full/10.1002/9780470974001.f205034. nez, J. Bisquert (Eds.), Photo[4] R.L. Doyle, M.E.G. Lyons, in: S. Gime electrochemical Solar Fuel Production: from Basic Principles to Advanced Devices, Springer International Publishing, Cham, 2016, pp. 41e104. [5] D.E. Hall, J. Electrochem. Soc. 132 (1985) 41C. [6] R.D.L. Smith, C. Pasquini, S. Loos, P. Chernev, K. Klingan, P. Kubella, M.R. Mohammadi, D. Gonzalez-Flores, H. Dau, Nat. Commun. 8 (2017) 2022. lez-Flores, K. Klingan, P. Chernev, S. Loos, M.R. Mohammadi, [7] D. Gonza C. Pasquini, P. Kubella, I. Zaharieva, R.D.L. Smith, H. Dau, Sustainable Energy Fuels 2 (2018) 1986. [8] A. Bergmann, E. Martinez-Moreno, D. Teschner, P. Chernev, M. Gliech, J.F. de Araújo, T. Reier, H. Dau, P. Strasser, Nat. Commun. 6 (2015) 8625. [9] L. Trotochaud, S.L. Young, J.K. Ranney, S.W. Boettcher, J. Am. Chem. Soc. 136 (2014) 6744. [10] S. Klaus, Y. Cai, M.W. Louie, L. Trotochaud, A.T. Bell, J. Phys. Chem. C 119 (2015) 7243. [11] M.W. Louie, A.T. Bell, J. Am. Chem. Soc. 135 (2013) 12329. [12] M.S. Burke, L.J. Enman, A.S. Batchellor, S. Zou, S.W. Boettcher, Chem. Mater. 27 (2015) 7549. [13] M.S. Burke, M.G. Kast, L. Trotochaud, A.M. Smith, S.W. Boettcher, J. Am. Chem. Soc. 137 (2015) 3638. [14] D.A. Corrigan, J. Electrochem. Soc. 134 (1987) 377. [15] C.C.L. McCrory, S. Jung, I.M. Ferrer, S.M. Chatman, J.C. Peters, T.F. Jaramillo, J. Am. Chem. Soc. 137 (2015) 4347. [16] C.C.L. McCrory, S. Jung, J.C. Peters, T.F. Jaramillo, J. Am. Chem. Soc. 135 (2013) 16977. [17] J.R. Swierk, S. Klaus, L. Trotochaud, A.T. Bell, T.D. Tilley, J. Phys. Chem. C 119 (2015) 19022. vot, R.D. Fagan, S. Trudel, C.P. Berlinguette, J. Am. Chem. [18] R.D.L. Smith, M.S. Pre Soc. 135 (2013) 11580. [19] L. Trotochaud, J.K. Ranney, K.N. Williams, S.W. Boettcher, J. Am. Chem. Soc. 134 (2012) 17253. € hring, B. Kieback, [20] T. Rauscher, C.I. Müller, A. Gabler, T. Gimpel, M. Ko €ntzsch, Electrochim. Acta 247 (2017) 1130. W. Schade, L. Ro € hring, B. Kieback, L. Ro € ntzsch, [21] A. Gabler, C.I. Müller, T. Rauscher, M. Ko W. Schade, Int. J. Hydrogen Energy 42 (2017) 10826. lez-Buch, I. Herraiz-Cardona, E. Ortega, J. García-Anto  n, V. Pe rez[22] C. Gonza Herranz, Int. J. Hydrogen Energy 38 (2013) 10157. [23] J. Balej, Int. J. Hydrogen Energy 10 (1985) 89. [24] A.J. Smith, D.L. Trimm, Annu. Rev. Mater. Res. 35 (2005) 127. [25] D.E. Hall, J. Electrochem. Soc. 128 (1981) 740. [26] S. Rausch, J. Electrochem. Soc. 143 (1996) 2852. €ntzsch, Int. J. Hydrogen [27] T. Rauscher, C.I. Bern€ acker, U. Mühle, B. Kieback, L. Ro Energy 44 (2019) 6392. [28] G. Kreysa, B. Håkansson, J. Electroanal. Chem. Interfacial Electrochem. 201 (1986) 61. [29] Z.A. Munir, D.V. Quach, M. Ohyanagi, J. Am. Ceram. Soc. 94 (2011) 1. [30] D. Giuntini, J. Raethel, M. Herrmann, A. Michaelis, E.A. Olevsky, D. Bouvard, J. Am. Ceram. Soc. 98 (2015) 3529. [31] M. Hakamada, Y. Yamada, T. Nomura, Y. Chen, H. Kusuda, M. Mabuchi, Mater. Trans. 46 (2005) 2624. [32] C.E. Wen, J.Y. Xiong, Y.C. Li, P.D. Hodgson, Phys. Scr. T139 (2010) 14070. [33] B. Arifvianto, J. Zhou, Materials 7 (2014) 3588. [34] A. Bansiddhi, D.C. Dunand, Intermetallics 15 (2007) 1612. [35] Y. Kanoko, K. Ameyama, S. Tanaka, B. Hefler, Powder Metall. 57 (2014) 168. [36] C.E. Wen, M. Mabuchi, Y. Yamada, K. Shimojima, H. Hosokawa, Y. Chino, T. Asahina, Mater. Sci. Forum 426e432 (2003) 417. [37] F. Zhang, E. Burkel, in: A. Laskovski (Ed.), Biomedical Engineering, Trends in Materials Science, InTech, 2011.

[38] D. Handtrack, F. Despang, C. Sauer, B. Kieback, N. Reinfried, Y. Grin, Mater. Sci. Eng. A 437 (2006) 423. € ntzsch, Int. J. Hydrogen [39] T. Rauscher, C.I. Müller, A. Schmidt, B. Kieback, L. Ro Energy 41 (2016) 2165. €ntzsch, [40] C.I. Müller, K. Sellschopp, M. Tegel, T. Rauscher, B. Kieback, L. Ro J. Power Sources 304 (2016) 196. [41] D.R. Lide (Ed.), CRC Handbook of Chemistry and Physics, CRC Press, Boca Raton Fla, 2005. [42] M. Tegel, DiffDiff (2013) in-house software: V 0.82. €rber, B. Kieback, J. Nucl. [43] I. Hilger, M. Tegel, M.J. Gorley, P.S. Grant, T. Weißga Mater. 447 (2014) 242.  , M. Alsabet, G. Jerkiewicz, ACS applied materials & interfaces 4 [44] M. Grden (2012) 3012.  ttir, M. Caspersen, S. Egelund, P. Møller, Electrochim. Acta 142 [45] C. Kjartansdo (2014) 324. [46] I.J. Godwin, M.E.G. Lyons, Electrochem. Commun. 32 (2013) 39. [47] H. Salmang, H. Scholze, R. Telle, Keramik, German Edition, Springer, Dordrecht, 2007. [48] W. Schatt, Pulvermetallurgie: Technologien und Werkstoffe, second ed., Springer-Verlag Berlin Heidelberg, Berlin, Heidelberg, 2007. [49] K. Ishizaki, S. Komarneni, M. Nanko, Porous Materials: Process Technology and Applications, Springer US, Boston, MA, s.l., 1998. [50] C. Hitz, A. Lasia, J. Electroanal. Chem. 500 (2001) 213. [51] L. Chen, J. Electrochem. Soc. 139 (1992) 3214. [52] L.M. Gassa, J.R. Vilche, M. Ebert, K. Jüttner, W.J. Lorenz, J. Appl. Electrochem. 20 (1990) 677. [53] A. Kellenberger, N. Vaszilcsin, W. Brandl, N. Duteanu, Int. J. Hydrogen Energy 32 (2007) 3258. [54] A. Kellenberger, N. Vaszilcsin, W. Brandl, J. Solid State Electrochem. 11 (2006) 84. [55] A. Lasia, A. Rami, J. Electroanal. Chem. Interfacial Electrochem. 294 (1990) 123. [56] P.W.T. Lu, S. Srinivasan, J. Electrochem. Soc. 125 (1978) 265. [57] S. Trasatti, O.A. Petrii, Pure Appl. Chem. 63 (1991) 711. [58] H. Vogt, R.J. Balzer, Electrochim. Acta 50 (2005) 2073.  La [59] B.M. Jovi c, U.C. cnjevac, V.D. Jovi c, N.V. Krstaji c, J. Electroanal. Chem. 754 (2015) 100. [60] D.S. Hall, C. Bock, B.R. MacDougall, J. Electrochem. Soc. 160 (2013) F235. [61] D.S. Hall, D.J. Lockwood, C. Bock, B.R. MacDougall, Proceedings. Mathematical, physical, and engineering sciences 471 (2015) 20140792. [62] M.E.G. Lyons, M.P. Brandon, Int. J. Electrochem. Sci. 3 (2008) 1386. [63] H. Bode, K. Dehmelt, J. Witte, Electrochim. Acta 11 (1966) 1079. [64] R. Barnard, C.F. Randell, J. Appl. Electrochem. 13 (1983) 27. [65] R. Barnard, C.F. Randell, J. Appl. Electrochem. 13 (1983) 89. [66] S.R. Mellsop, A. Gardiner, B. Johannessen, A.T. Marshall, Electrochim. Acta 168 (2015) 356. [67] R.L. Doyle, I.J. Godwin, M.P. Brandon, M.E.G. Lyons, Phys. Chem. Chem. Phys. 15 (2013) 13737. [68] B.S. Yeo, A.T. Bell, J. Phys. Chem. C 116 (2012) 8394. [69] L. Vra car, B.E. Conway, Electrochim. Acta 35 (1990) 1919. [70] P.W.T. Lu, S. Srinivasan, J. Electrochem. Soc. 125 (1978) 1416. [71] C. Bocca, Int. J. Hydrogen Energy 23 (1998) 247. [72] R.D.L. Smith, R.S. Sherbo, K.E. Dettelbach, C.P. Berlinguette, Chem. Mater. 28 (2016) 5635. [73] L. Trotochaud, S.W. Boettcher, Scripta Mater. 74 (2014) 25. [74] J.-Y. Huot, J. Electrochem. Soc. 136 (1989) 1933. [75] B.E. Conway, L. Bai, J. Chem. Soc., Faraday Trans. 1 (81) (1985) 1841. [76] B.E. Conway, L. Bai, Int. J. Hydrogen Energy 11 (1986) 533. [77] D.A. Harrington, B.E. Conway, Electrochim. Acta 32 (1987) 1703. [78] E. Gileadi, Physical Electrochemistry: Fundamentals, Techniques and Applications, Wiley-VCH-Verl., Weinheim, 2011. [79] M. Gong, Y. Li, H. Wang, Y. Liang, J.Z. Wu, J. Zhou, J. Wang, T. Regier, F. Wei, H. Dai, J. Am. Chem. Soc. 135 (2013) 8452. [80] F. Song, X. Hu, Nat. Commun. 5 (2014) 4477. €rlin, P. Chernev, J. Ferreira de Araújo, T. Reier, S. Dresp, B. Paul, [81] M. Go €hnert, H. Dau, P. Strasser, J. Am. Chem. Soc. 138 (2016) 5603. R. Kra €blinger, [82] K. Fominykh, P. Chernev, I. Zaharieva, J. Sicklinger, G. Stefanic, M. Do €cklein, C. Scheu, T. Bein, D. Fattakhova-Rohlfing, A. Müller, A. Pokharel, S. Bo ACS Nano 9 (2015) 5180. [83] C. Roy, B. Sebok, S.B. Scott, E.M. Fiordaliso, J.E. Sørensen, A. Bodin, D.B. Trimarco, C.D. Damsgaard, P.C.K. Vesborg, O. Hansen, I.E.L. Stephens, J. Kibsgaard, I. Chorkendorff, Nat Catal 1 (2018) 820. [84] G.N. Martelli, R. Ornelas, G. Faita, Electrochim. Acta 39 (1994) 1551. [85] A.C.C. Tseung, S. Jasem, Electrochim. Acta 22 (1977) 31. [86] S. Tanaka, Int. J. Hydrogen Energy 25 (2000) 481. [87] M. Okido, J. Electrochem. Soc. 140 (1993) 127. [88] R.K. Shervedani, A. Lasia, J. Appl. Electrochem. 29 (1999) 979.