**1. Introduction**

The manufacture of most carbon materials that are technologically relevant has strongly relied on the fossil fuel industry. The predicted gradual phase-out of this industry is thus expected to have a massive impact on the way technological carbons are made today, with carbon materials corporations being compelled to search for low-carbon alternatives, including precursors. In this scenario, biomass is expected to play a primary role as a precursor in several carbon markets in which its presence has been so far residual or non-existent [1–3].

Biomass is already a significant contributor to the industrial production of porous carbon materials. Activated carbon companies nowadays consume lignocellulosic biomass (wood and coconut shell mainly) to obtain porous carbon powders or granules, which are daily used products in a wide variety of activities including environmental and energy storage applications [4–6]. However, in a number of processes, 3D structured porous carbons (carbon monoliths) offer several advantages over powdered or granular ones [7–9]. Monolithic structures incorporate transport pores to reduce the pressure drop in continuous operations such as industrial filters [10,11]. They are also very relevant materials for heterogeneous catalytic reactors at the industrial level [12,13], and gas storage applications [14–16]. In all these applications, not only the porosity of the carbon material is

**Citation:** Llamas-Unzueta, R.; Ramírez-Montoya, L.A.; Menéndez, J.A.; Montes-Morán, M.A. Customised Microporous Carbon 3D Structures with Good Mechanical Properties and High Nitrogen Content Obtained from Whey Powders. *C* **2023**, *9*, 100. https:// doi.org/10.3390/c9040100

Academic Editor: Jandro L. Abot

Received: 29 August 2023 Revised: 27 September 2023 Accepted: 17 October 2023 Published: 24 October 2023

**Copyright:** © 2023 by the authors. Licensee MDPI, Basel, Switzerland. This article is an open access article distributed under the terms and conditions of the Creative Commons Attribution (CC BY) license (https:// creativecommons.org/licenses/by/ 4.0/).

of paramount importance but also how the 3D carbon structure endures the operating conditions in industrial plants. In other words, 3D structured porous carbons should also have good mechanical properties, which makes the current fossil fuel-derived resins and binders (i.e., pitch) imperative in the manufacture of carbon monoliths [17–19].

There have been significant efforts to cut off ties with this undesirable dependence. Carboxymethyl cellulose (CMC) has been used as a green binder with remarkable results in terms of the mechanical properties of the resulting activated carbon monoliths [20]. However, the use of binders to agglomerate activated carbon powders has traditionally important drawbacks related to pore clogging [21,22]. Strategies based on binder-free monoliths from biomass precursors are scarce. Indeed, a very interesting methodology comprises the formation of slurries from activating agents such as H3PO<sup>4</sup> or ZnCl<sup>2</sup> and lignocellulosic precursors [23–25]. These slurries are prone to be either uniaxially compressed or extruded to obtain monoliths that withstand carbonisation and washing of the activation by-products. On the other hand, sustainable approaches to a greener synthesis of 3D structured porous carbons comprise studies with a number of novel biomass-derived precursors including polysaccharides [26], tannins [27], or novel biomass wastes (i.e., beyond lignocellulosic precursors) [28]. Hydrothermal carbonisation (HTC) and other approaches also looked forward to obtaining 3D structured porous carbons, mainly carbon aerogels [29,30]. However, results in terms of the mechanical integrity of those green carbon monoliths have been, in general, rather limited.

We have recently discovered the possibility of obtaining 3D porous carbons with outstanding mechanical properties by direct carbonisation of whey powders poured on moulds [31]. Whey powders result after the spray-drying of liquid whey, a byproduct of cheese and casein production found in massive quantities in the dairy industry [32]. The relatively high organic matter (mainly lactose) load of liquid whey makes whey management a serious burden. Big dairy companies have invested important resources to valorise whey and an estimated 50% of liquid whey ends nowadays in food and feed industries [33,34]. Still, ca. 100 million tons of liquid whey are wasted each year worldwide [35]. Since, as detailed in [31], the mechanical properties of the 3D whey derived carbons compete with those prepared from conventional resins, whey should be then considered a robust biomass precursor to manufacture green carbon monoliths that can be used, for example, as bone scaffolds [36].

The porosity of the 3D whey-derived carbons have been also characterised in previous studies. The carbonisation of whey powders renders materials that are mainly macroporous, with limited meso- and micropore development, as determined by N<sup>2</sup> adsorption at cryogenic temperatures [31]. The aim of this work is to study the activation of whey carbon monoliths to improve their micro- and mesoporosity, which is compulsory for their use as adsorbents or catalysts/catalyst supports. Although KOH activation of whey powders has been reported before [37], as it has been that of non-fat powdered milk, a precursor similar to whey [38], the final product of these studies are activated carbon powders, which is not the purpose of this work.

#### **2. Materials and Methods**

#### *2.1. Synthesis and Activation of the Whey-Derived Carbon Monoliths*

The preparation of the custom-made microporous carbon 3D structures from whey comprises two steps: (i) pre-conformation of the 3D structure; (ii) carbonisation and/or activation.

For the pre-conformation of the 3D structures, whey powder (W) (CAPSA-Food, Granda, Spain) was poured into a 12 mm inner diameter glass tube. Details of W including its chemical composition have been reported elsewhere [31]. The pre-conformation step was carried out at 150 ◦C for 1 h in a lab stove. After this step, the W powder transformed into a solid rod of 12 mm diameter. The rod was then cut in cylinders of 30 mm height using a circular saw.

The cylinders were further carbonised at different temperatures in a tubular furnace (Carbolite, Hope Valley, UK) under N<sup>2</sup> (100 mL/min) with a heating rate and dwell time

of 10 ◦C/min and 1.5 h, respectively. The carbonised samples were labelled as CW\_X, with X being the carbonisation temperature. Some of the carbonised whey cylinders were demineralised (samples CW\_X\_D) by acid (HCl) plus water washing in an ultrasonic bath.

Thermally activated samples (TAW) were first heat treated at up to 850 ◦C for 1.5 h under 100 mL/min N2. After carbonisation, the temperature of the oven was set to the activation temperature (800–850 ◦C), the N<sup>2</sup> flux was switched to 50 mL/min CO<sup>2</sup> and kept at that activation temperature for 1–3 h. Finally, the monolith was cooled down in N<sup>2</sup> (100 mL/min). The TAW\_X\_Y samples' labelling code includes the activation temperature (X) followed by the dwelling time (Y).

Chemically activated samples (CASs) were prepared using both KOH and H3PO<sup>4</sup> as activating agents (CH). KOH powder, H3PO<sup>4</sup> (85% in water) and fuming HCl were purchased from Sigma-Aldrich (Schnelldorf, Germany). Milli-Q water was used for washing and impregnation procedures. The direct chemical activation of the pre-conformed monoliths was not feasible and a thermal stabilisation step was required prior activation. In the case of H3PO<sup>4</sup> activation, the pre-conformed cylinders were heated up to 250 ◦C in the tubular furnace under N<sup>2</sup> (100 mL/min) for 1.5 h, whereas for KOH activation, the stabilisation temperature was set at 450 ◦C. The thermally stabilised monoliths (S) were then impregnated with different proportions of either KOH or H3PO4. This impregnation step was carried out by sonicating the monolith in 100 mL of a water solution of the CH at 80 ◦C for 2 h and then dried at 100 ◦C. It is important to mention that the integrity of the monolith remained unaltered after this stage. Afterwards, the impregnated monolith was placed in an alumina crucible and thermally treated at different temperatures for 1.5 h with a heating ramp of 10 ◦C/min under N<sup>2</sup> atmosphere (100 mL/min). After cooling down in N2, the CAS samples obtained using KOH were washed with 1 M HCl under US during three cycles of 15 min each followed by vacuum filtration for each batch. Then, samples were washed with distilled water in an ultrasonic bath and filtrated. This procedure was performed for three consecutive 15 min washing cycles. The samples were finally dried at 110 ◦C in a convection oven overnight. For phosphoric activated samples, five washing cycles were performed using only sonicated water. The resulting monoliths were labelled as CAS\_X\_CH\_Z, X being the heat treatment temperature, CH the activating agent (K for KOH and P for H3PO4) and Z the S:CH impregnation ratio.

### *2.2. Characterisation Methods*

The thermogravimetric (TG) analyses of the W powders were carried out in a TGA Q5000 device from TA Instruments (New Castle, DE, USA) under N<sup>2</sup> and CO<sup>2</sup> atmospheres (20 mL/min) using a heating rate of 10 ◦C/min from room temperature to 1000 ◦C with sample masses of ca. 20 mg. For ash and moisture contents of the carbonised and activated monoliths (also ca. 20 mg), TGs were obtained using 20 mL/min of N<sup>2</sup> up to 900 ◦C and then switching to air (20 mL/min) for 30 min at that temperature. The moisture contents were calculated from the loss of weight at 110 ◦C. The ash content was estimated by the difference of the total loss of weight of the sample after combustion at 900 ◦C on air. The elemental analysis of the carbon monoliths was carried out in LECO (Geleen, The Netherlands) apparatuses (LECO CHNS-932, and LECO VTF-900 for oxygen content determination). The point of zero charge of the monoliths (pHPZC) was determined by mass titration [39].

Nitrogen adsorption–desorption isotherms at −196 ◦C were performed in a Micromeritics Tristar II volumetric adsorption system, after outgassing the samples at 120 ◦C overnight. The specific surface area (SBET) and micropore volume (Vmicro) were calculated using the Brunauer–Emmett–Teller (BET) and Dubinin–Radushkevich (DR) models, respectively, while the pore size distributions (PSD) were calculated using the Density Functional Theory (DFT) model. The helium densities (ρHe) were determined in an Accu-Pyc 1330 pycnometer (Micromeritics, Norcross, GA, USA). For the Hg intrusion technique, an AutoPore IV porosimeter (Micromeritics, GA, USA) was used. The Hg densities (ρHg) and corresponding PSDs were determined from the Hg intrusion data up to a maximum

operating pressure of 227 MPa. For both He and Hg densities determination, samples were outgassed at 120 ◦C overnight. outgassed at 120 °C overnight. Compressive strengths (Su) were evaluated in an Instron Model 8562 (Instron, MA, USA) device with a cell load of 10 kN and a test velocity of 0.5 mm/min. The final di-

cuPyc 1330 pycnometer (Micromeritics, GA, USA). For the Hg intrusion technique, an AutoPore IV porosimeter (Micromeritics, GA, USA) was used. The Hg densities (ρHg) and corresponding PSDs were determined from the Hg intrusion data up to a maximum operating pressure of 227 MPa. For both He and Hg densities determination, samples were

Compressive strengths (Su) were evaluated in an Instron Model 8562 (Instron, Norwood, MA, USA) device with a cell load of 10 kN and a test velocity of 0.5 mm/min. The final dimensions of the tested specimens were 12.25 ± 0.2 mm diameter and 22.5 ± 0.2 mm height. The compressive strength value is the maximum stress (load/area) that the specimen stood during the test. The elastic compressive moduli (E) were obtained from the slope of the linear elastic region of the stress vs. strain curve. For a given sample, the S<sup>u</sup> and E values reported are the average of three compressive tests. mensions of the tested specimens were 12.25 ± 0.2 mm diameter and 22.5 ± 0.2 mm height. The compressive strength value is the maximum stress (load/area) that the specimen stood during the test. The elastic compressive moduli (E) were obtained from the slope of the linear elastic region of the stress vs. strain curve. For a given sample, the Su and E values reported are the average of three compressive tests. **3. Results and Discussion** 

#### **3. Results and Discussion** *3.1. Customised 3D Structures Made of Carbonised Whey Powders*

#### *3.1. Customised 3D Structures Made of Carbonised Whey Powders* As shown in Figure 1, 3D monolithic porous carbons with different shapes and sizes

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As shown in Figure 1, 3D monolithic porous carbons with different shapes and sizes were obtained by the simple carbonisation of pre-conformed W powders, depending on the selected mould (cubes, cylinders, spheres, hexagons, etc.). Even more sophisticated geometries, such as nuts or hollow cylinders, are possible. Also, the mechanical stability of the whey-derived carbon monoliths bears machining processes to obtain customised materials for a variety of applications. As pointed out in previous works [31], whey particles stick to each other by the self-reaction between the precursor constituents (lactose and whey proteins) during the heat treatment. Whey particles do not melt with temperature and both the pre-conformation and carbonisation/activation of the 3D structures is carried out in absence of any external template, binders or pressure. It should be finally mentioned that quantities of inorganic matter (6–13% ash contents, depending on the W batch and/or brand) remain in the carbon monoliths, mainly composed of Ca, K, Na, and P [31,40]. were obtained by the simple carbonisation of pre-conformed W powders, depending on the selected mould (cubes, cylinders, spheres, hexagons, etc.). Even more sophisticated geometries, such as nuts or hollow cylinders, are possible. Also, the mechanical stability of the whey-derived carbon monoliths bears machining processes to obtain customised materials for a variety of applications. As pointed out in previous works [31], whey particles stick to each other by the self-reaction between the precursor constituents (lactose and whey proteins) during the heat treatment. Whey particles do not melt with temperature and both the pre-conformation and carbonisation/activation of the 3D structures is carried out in absence of any external template, binders or pressure. It should be finally mentioned that quantities of inorganic matter (6–13% ash contents, depending on the W batch and/or brand) remain in the carbon monoliths, mainly composed of Ca, K, Na, and P [31,40].

**Figure 1.** Examples of customised whey-derived carbon 3D structures made by simple moulding, or by moulding + machining. **Figure 1.** Examples of customised whey-derived carbon 3D structures made by simple moulding, or by moulding + machining.

#### *3.2. Thermal Activation of the Whey-Derived 3D Carbons 3.2. Thermal Activation of the Whey-Derived 3D Carbons*

As mentioned in the Introduction, whey carbon monoliths are essentially macroporous materials (maximum of the Hg intrusion PSD centred at 20–40 µm, see Section 3.4) and show no significant micro or mesoporosity, as measured by N2 adsorption at −196 °C (Figure 2a). However, the adsorption of CO2 at 0 °C on the same materials (Figure 2b) is significant and reveals the ultra-microporosity (pore sizes below 0.7 nm) of these materials. The ultra-microporosity is incipient in the low-temperature carbonised As mentioned in the Introduction, whey carbon monoliths are essentially macroporous materials (maximum of the Hg intrusion PSD centred at 20–40 µm, see Section 3.4) and show no significant micro or mesoporosity, as measured by N<sup>2</sup> adsorption at −196 ◦C (Figure 2a). However, the adsorption of CO<sup>2</sup> at 0 ◦C on the same materials (Figure 2b) is significant and reveals the ultra-microporosity (pore sizes below 0.7 nm) of these materials. The ultra-microporosity is incipient in the low-temperature carbonised monoliths CW\_450 and increases substantially as the carbonisation temperature rises up to 750 and 850 ◦C. Finally, higher carbonisation temperatures (CW\_1000) result in a decrease in the volume of CO<sup>2</sup> adsorbed, thus suggesting the collapse of the narrow micropores. The relevant CO<sup>2</sup> adsorption on these materials, especially CW\_750 and CW\_850, opens the possibility of

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transforming the whey-derived carbon 3D structures into activated carbon monoliths by thermal activation. CW\_850, opens the possibility of transforming the whey-derived carbon 3D structures into activated carbon monoliths by thermal activation.

monoliths CW\_450 and increases substantially as the carbonisation temperature rises up to 750 and 850 °C. Finally, higher carbonisation temperatures (CW\_1000) result in a decrease in the volume of CO2 adsorbed, thus suggesting the collapse of the narrow micropores. The relevant CO2 adsorption on these materials, especially CW\_750 and

**Figure 2.** (**a**) N2 adsorption isotherms (−196 °C); and (**b**) CO2 adsorption isotherms (0 °C) on whey monoliths carbonised at different temperatures. **Figure 2.** (**a**) N<sup>2</sup> adsorption isotherms (−196 ◦C); and (**b**) CO<sup>2</sup> adsorption isotherms (0 ◦C) on whey monoliths carbonised at different temperatures.

This was further confirmed by the TG analysis of the precursor W. Figure 3 shows the TG and DTG profiles of W heated in both N2 and CO2 atmospheres. Discussing first the TG profile of whey carbonisation, the TG curve is characterised by five stages. The first stage with the DTG peak at ca. 80 °C is attributed to the evaporation of the moisture remaining in the W sample. The second stage (130–160 °C) corresponds to the loss of crystalline water present in the α-lactose monohydrate (spray dried whey particles are a matrix of amorphous lactose in which lactose monohydrate crystals, proteins, fats and minerals are embedded). The reactions occurring in this temperature range are critical for the formation of 3D structures from W that withstand carbonisation. The release of crystalline water in this second stage is expected to boost the formation of melanoidins through the Maillard reactions between the lactose and the proteins in W. The successive thermal events (from 160 °C onwards) are thus complex and comprise the thermal evolution of unreacted lactose and whey proteins, on one side [31], and melanoidins, on the This was further confirmed by the TG analysis of the precursor W. Figure 3 shows the TG and DTG profiles of W heated in both N<sup>2</sup> and CO<sup>2</sup> atmospheres. Discussing first the TG profile of whey carbonisation, the TG curve is characterised by five stages. The first stage with the DTG peak at ca. 80 ◦C is attributed to the evaporation of the moisture remaining in the W sample. The second stage (130–160 ◦C) corresponds to the loss of crystalline water present in the α-lactose monohydrate (spray dried whey particles are a matrix of amorphous lactose in which lactose monohydrate crystals, proteins, fats and minerals are embedded). The reactions occurring in this temperature range are critical for the formation of 3D structures from W that withstand carbonisation. The release of crystalline water in this second stage is expected to boost the formation of melanoidins through the Maillard reactions between the lactose and the proteins in W. The successive thermal events (from 160 ◦C onwards) are thus complex and comprise the thermal evolution of unreacted lactose and whey proteins, on one side [31], and melanoidins, on the other [41,42]. *C* **2023**, *9*, x FOR PEER REVIEW 6 of 18

**Figure 3.** TG and DTG of the whey powders under N2 and CO2 atmospheres. **Figure 3.** TG and DTG of the whey powders under N<sup>2</sup> and CO<sup>2</sup> atmospheres.

The TG of W in CO2 is, as expected, virtually identical to that in N2 during most of the temperature scan. Gasification starts at 720 °C, and the quick mass drop, also reflected in the sharpness of the corresponding CO2 DTG band with maximum at 935 °C, indicates

The activation of the W chars (CW\_850) was thus carried out at 800 °C and 850 °C. Before showing and discussing the results obtained, a fundamental comment regarding the thermal activation of these monoliths should be pointed out. The thermal activation of the stocky 3D structures of carbonised whey in the tubular oven was quite tricky. Considering that, even when working with powders, the standard procedures of CO2 activation recommend the deposit of thin layers over the crucible in order to facilitate the contact between the gas and the carbon material, the activation of a monolith represents a challenge. The CO2 needs to diffuse into the whole cylinder in order to obtain a homogeneous activation, i.e., to minimise the differences of activation between the skin and the core of the monolith [44]. N2 isotherms (−196 °C) of the monolith skin and core were thus measured by removing the outermost (approx. 1 mm) crust of the cylinders after activation. Although we were confident that the open macroporosity of the whey-derived chars would ease the diffusion of the CO2 inside the cylinders, the initial experiments failed to provide homogeneous activation of the pieces. Regardless of the changes in the composition (relative proportion of CO2 and N2) and flux of the inlet gas, those initial experiments brought about either non-activated monoliths or only skin-activated monoliths. It was necessary to plug a substantial amount of quartz wool at the tube inlet (approx. 10 cm length), that helped to spread the laminar flow of CO2, to succeed with differences between the skin and core being less than 10%. The results shown in Figure 4 and Table 1

biomass chars such as coconut char [43]. This high reactivity of the whey char is a consequence of two contributions: (i) the high ultra-micropore volume of the CWs obtained at carbonisation temperatures above 700 °C (Figure 2b); and (ii) the existence of mineral matter catalysing the gasification. Analysis of the TG/DTG CO2 profile of W suggests that the best temperature range for thermal activation of W chars is 800–850 °C, temperatures

at which the gasification would be kept under (relative) control.

correspond to the measurements of the whole (skin + core) monolith.

The TG of W in CO<sup>2</sup> is, as expected, virtually identical to that in N<sup>2</sup> during most of the temperature scan. Gasification starts at 720 ◦C, and the quick mass drop, also reflected in the sharpness of the corresponding CO<sup>2</sup> DTG band with maximum at 935 ◦C, indicates the relatively high reactivity of the W char when compared, for example, to conventional biomass chars such as coconut char [43]. This high reactivity of the whey char is a consequence of two contributions: (i) the high ultra-micropore volume of the CWs obtained at carbonisation temperatures above 700 ◦C (Figure 2b); and (ii) the existence of mineral matter catalysing the gasification. Analysis of the TG/DTG CO<sup>2</sup> profile of W suggests that the best temperature range for thermal activation of W chars is 800–850 ◦C, temperatures at which the gasification would be kept under (relative) control.

The activation of the W chars (CW\_850) was thus carried out at 800 ◦C and 850 ◦C. Before showing and discussing the results obtained, a fundamental comment regarding the thermal activation of these monoliths should be pointed out. The thermal activation of the stocky 3D structures of carbonised whey in the tubular oven was quite tricky. Considering that, even when working with powders, the standard procedures of CO<sup>2</sup> activation recommend the deposit of thin layers over the crucible in order to facilitate the contact between the gas and the carbon material, the activation of a monolith represents a challenge. The CO<sup>2</sup> needs to diffuse into the whole cylinder in order to obtain a homogeneous activation, i.e., to minimise the differences of activation between the skin and the core of the monolith [44]. N<sup>2</sup> isotherms (−196 ◦C) of the monolith skin and core were thus measured by removing the outermost (approx. 1 mm) crust of the cylinders after activation. Although we were confident that the open macroporosity of the whey-derived chars would ease the diffusion of the CO<sup>2</sup> inside the cylinders, the initial experiments failed to provide homogeneous activation of the pieces. Regardless of the changes in the composition (relative proportion of CO<sup>2</sup> and N2) and flux of the inlet gas, those initial experiments brought about either non-activated monoliths or only skin-activated monoliths. It was necessary to plug a substantial amount of quartz wool at the tube inlet (approx. 10 cm length), that helped to spread the laminar flow of CO2, to succeed with differences between the skin and core being less than 10%. The results shown in Figure 4 and Table 1 correspond to the measurements of the whole (skin + core) monolith. *C* **2023**, *9*, x FOR PEER REVIEW 7 of 18

**Figure 4.** (**a**) N2 adsorption isotherms (−196 °C); and (**b**) DFT pore size distributions of the thermally (CO2) activated monoliths derived from whey. **Figure 4.** (**a**) N<sup>2</sup> adsorption isotherms (−196 ◦C); and (**b**) DFT pore size distributions of the thermally (CO<sup>2</sup> ) activated monoliths derived from whey.

**Vmeso <sup>c</sup> (cm3/g)** 

**ρHe (g/cm3)**  **Yield d (%)** 

Vmeso = Vtotal − Vmicro; d expressed as

**Burn-Off <sup>e</sup> (%)** 

**Table 1.** Selected textural parameters and yields of the thermally activated carbon monoliths.

**Vmicro <sup>b</sup> (cm3/g)** 

CW\_850 <10 0.001 - - 2.00 22.4 0 TAW\_800\_1 500 0.287 0.205 0.082 2.07 17.2 23.2

TAW\_800\_3 834 0.529 0.332 0.197 1.98 12.3 45.1 TAW\_850\_1 839 0.516 0.331 0.185 2.01 12.3 45.1 TAW\_850\_1.5 981 0.689 0.377 0.312 2.01 10.1 54.9 TAW\_850\_2 1140 0.995 0.442 0.553 1.99 7.1 68.3

Thermal activation is effective at both temperatures, rendering materials that combine micro and mesoporosity (isotherm Type IV, Figure 4a). The monoliths activated at 800 °C show a limited development of the microporosity, with SBET maximum values of 840 m2/g after 3 h of dwell time. This very same value is obtained after only 1 h of CO2 activation at 850 °C, whereas 2 h activation at this higher temperature brings about materials with SBET values well above 1100 m2/g. Due to the high reactivity of the whey-derived carbon, this relatively high surface area is attained at the expense of a high burn-off (68.3%, Table 1) of the whey char, which strongly limits the mechanical integrity of the resulting 3D structure (Figure 5). Profusion of cracks in the whey monoliths makes the activation conditions (2 h at 850 °C) unfeasible for their practical use, thus establishing the temperature and dwell time boundaries for thermal activation of these 3D structures with CO2. On the other hand, looking at the textural parameters of the activated monoliths at 850 °C (Table 1), there is a considerable gap between the specific surface area of the TAW\_850\_1 and TAW\_850\_2. We decided then to perform an additional activation experiment with a 1.5 h dwell time (sample TAW\_850\_1.5); the results are also included in Figure 4 and Table 1. The TAW\_850\_1.5 monoliths were free of external crevices.

a Calculated at p/p0 = 0.99; b calculated with the DR method; c

(massmonolith/masswhey) × 100; e expressed as 100 − [(yieldCW\_850/yieldmonolith) × 100].

**Vtotal <sup>a</sup> (cm3/g)** 

**Sample SBET**

**(m2/g)** 


**Table 1.** Selected textural parameters and yields of the thermally activated carbon monoliths.

<sup>a</sup> Calculated at p/p<sup>0</sup> = 0.99; <sup>b</sup> calculated with the DR method; <sup>c</sup> <sup>V</sup>meso = Vtotal <sup>−</sup> <sup>V</sup>micro; d expressed as (massmonolith/masswhey) <sup>×</sup> 100; <sup>e</sup> expressed as 100 − [(yieldCW\_850/yieldmonolith) × 100].

Thermal activation is effective at both temperatures, rendering materials that combine micro and mesoporosity (isotherm Type IV, Figure 4a). The monoliths activated at 800 ◦C show a limited development of the microporosity, with SBET maximum values of 840 m2/g after 3 h of dwell time. This very same value is obtained after only 1 h of CO<sup>2</sup> activation at 850 ◦C, whereas 2 h activation at this higher temperature brings about materials with SBET values well above 1100 m2/g. Due to the high reactivity of the whey-derived carbon, this relatively high surface area is attained at the expense of a high burn-off (68.3%, Table 1) of the whey char, which strongly limits the mechanical integrity of the resulting 3D structure (Figure 5). Profusion of cracks in the whey monoliths makes the activation conditions (2 h at 850 ◦C) unfeasible for their practical use, thus establishing the temperature and dwell time boundaries for thermal activation of these 3D structures with CO2. On the other hand, looking at the textural parameters of the activated monoliths at 850 ◦C (Table 1), there is a considerable gap between the specific surface area of the TAW\_850\_1 and TAW\_850\_2. We decided then to perform an additional activation experiment with a 1.5 h dwell time (sample TAW\_850\_1.5); the results are also included in Figure 4 and Table 1. The TAW\_850\_1.5 monoliths were free of external crevices. *C* **2023**, *9*, x FOR PEER REVIEW 8 of 18

**Figure 5.** An example of a monolith with crevices due to overactivation. This picture corresponds to TAW\_850\_2, but it is also representative of the chemically activated monoliths. **Figure 5.** An example of a monolith with crevices due to overactivation. This picture corresponds to TAW\_850\_2, but it is also representative of the chemically activated monoliths.

The gasification of the whey monoliths under the tested conditions encompasses not only the formation of micropores but also an enhancement of the mesoporosity. The pore size distribution in the mesopore range of these materials is very similar for all materials except TAW\_850\_1.5 (Figure 4b). They are wide distributions in which several maxima can be spotted at ca. 20, 35, and 50 nm. In the case of the TAW\_850\_1.5 mesopore size distribution, the maximum at 20 nm seems to further split into two bands at 10–12 nm The gasification of the whey monoliths under the tested conditions encompasses not only the formation of micropores but also an enhancement of the mesoporosity. The pore size distribution in the mesopore range of these materials is very similar for all materials except TAW\_850\_1.5 (Figure 4b). They are wide distributions in which several maxima can be spotted at ca. 20, 35, and 50 nm. In the case of the TAW\_850\_1.5 mesopore size distribution, the maximum at 20 nm seems to further split into two bands at 10–12 nm

and 25 nm. The evolution of the micropore and mesopore volumes (Table 1) as the activation conditions become harsher confirms that a 850 °C and 2 h dwell is a point of in-

**Figure 6.** Specific surface area (SBET) and micropore volume (Vmicro) as a function of the thermal activation burn-off of the whey-derived activated carbon monoliths. Void and solid symbols corre-

0.20

0.25

0.30

Vmicro (cm3/g)

0.35

0.40

0.45

Figure 6 shows the relation between the burn-off and the specific surface area and micropore volume of the thermally activated samples. A good linear correlation exists in

= 0.996)

=0.996)

spond to the TAW\_800 and TAW\_850 monoliths, respectively.

20 30 40 50 60 70

Burn-off (%)

*3.3. Chemical Activation of the Whey-Derived 3D Structures* 

micropore network.

, SBET (r<sup>2</sup>

, Vmicro (r<sup>2</sup>

both cases.

400

500

600

700

800

SBET (m2/g)

900

1000

1100

1200

and 25 nm. The evolution of the micropore and mesopore volumes (Table 1) as the activation conditions become harsher confirms that a 850 ◦C and 2 h dwell is a point of inflection with Vmeso > Vmicro, thus suggesting that gasification of the char is widening the micropore network. and 25 nm. The evolution of the micropore and mesopore volumes (Table 1) as the activation conditions become harsher confirms that a 850 °C and 2 h dwell is a point of inflection with Vmeso > Vmicro, thus suggesting that gasification of the char is widening the micropore network.

**Figure 5.** An example of a monolith with crevices due to overactivation. This picture corresponds to

The gasification of the whey monoliths under the tested conditions encompasses not only the formation of micropores but also an enhancement of the mesoporosity. The pore size distribution in the mesopore range of these materials is very similar for all materials except TAW\_850\_1.5 (Figure 4b). They are wide distributions in which several maxima can be spotted at ca. 20, 35, and 50 nm. In the case of the TAW\_850\_1.5 mesopore size distribution, the maximum at 20 nm seems to further split into two bands at 10–12 nm

TAW\_850\_2, but it is also representative of the chemically activated monoliths.

Figure 6 shows the relation between the burn-off and the specific surface area and micropore volume of the thermally activated samples. A good linear correlation exists in both cases. Figure 6 shows the relation between the burn-off and the specific surface area and micropore volume of the thermally activated samples. A good linear correlation exists in both cases.

*C* **2023**, *9*, x FOR PEER REVIEW 8 of 18

**Figure 6.** Specific surface area (SBET) and micropore volume (Vmicro) as a function of the thermal activation burn-off of the whey-derived activated carbon monoliths. Void and solid symbols correspond to the TAW\_800 and TAW\_850 monoliths, respectively. **Figure 6.** Specific surface area (SBET) and micropore volume (Vmicro) as a function of the thermal activation burn-off of the whey-derived activated carbon monoliths. Void and solid symbols correspond to the TAW\_800 and TAW\_850 monoliths, respectively.

#### *3.3. Chemical Activation of the Whey-Derived 3D Structures 3.3. Chemical Activation of the Whey-Derived 3D Structures*

KOH and H3PO<sup>4</sup> were selected as activating agents (CH) due to their well-known capacity to generate a microporous network into the carbon matrix. As detailed in the Section 2, the chemical activation was carried out over thermally stabilised monoliths. The direct activation of the whey brought about melted structures that did not keep the shape of the mould containing them, obtaining either powders (KOH activation) or foamy structures (H3PO<sup>4</sup> activation). On the other hand, the issue that was pointed out when discussing the results of the thermal activation, regarding the homogeneity of the activated materials, was much less critical in the case of the chemical activation of the thermally stabilised monoliths. The macroporosity already present in these stabilised monoliths was well interconnected and the CH impregnation step was effective in reaching the inner parts of the stabilised monoliths. In addition, in the case of KOH it is also expected that the melting of the salt at temperatures above 360 ◦C will help the infiltration of CH into the monoliths and thus lead to a more homogeneous activation.

#### 3.3.1. KOH Activation

Figure 7 shows the N<sup>2</sup> adsorption isotherms (−196 ◦C) of the KOH-activated samples and Table 2 summarises the principal textural parameters calculated from those isotherms. All N<sup>2</sup> isotherms are Type I with well-defined plateaus corresponding to highly microporous materials with slit-shaped pores in the micropore range, with the exception of CAS\_750\_K\_1:1, which is a mixed Type I and IV isotherm with a hysteresis loop denoting the presence of mesopores (Vmeso = 0.137 cm3/g, Table 2).

noting the presence of mesopores (Vmeso = 0.137 cm3/g, Table 2).

3.3.1. KOH Activation

**Figure 7.** (**a**) N2 adsorption isotherms (−196 °C); and (**b**) DFT pore size distributions of the chemically (KOH) activated monoliths derived from whey. Inset of (**b**) is a zoom of a selected area (5–50 nm). **Figure 7.** (**a**) N<sup>2</sup> adsorption isotherms (−196 ◦C); and (**b**) DFT pore size distributions of the chemically (KOH) activated monoliths derived from whey. Inset of (**b**) is a zoom of a selected area (5–50 nm).

KOH and H3PO4 were selected as activating agents (CH) due to their well-known capacity to generate a microporous network into the carbon matrix. As detailed in the Materials and Methods section, the chemical activation was carried out over thermally stabilised monoliths. The direct activation of the whey brought about melted structures that did not keep the shape of the mould containing them, obtaining either powders (KOH activation) or foamy structures (H3PO4 activation). On the other hand, the issue that was pointed out when discussing the results of the thermal activation, regarding the homogeneity of the activated materials, was much less critical in the case of the chemical activation of the thermally stabilised monoliths. The macroporosity already present in these stabilised monoliths was well interconnected and the CH impregnation step was effective in reaching the inner parts of the stabilised monoliths. In addition, in the case of KOH it is also expected that the melting of the salt at temperatures above 360 °C will help the infiltration of CH into the monoliths and thus lead to a more homogeneous activation.

Figure 7 shows the N2 adsorption isotherms (−196 °C) of the KOH-activated samples and Table 2 summarises the principal textural parameters calculated from those isotherms. All N2 isotherms are Type I with well-defined plateaus corresponding to highly microporous materials with slit-shaped pores in the micropore range, with the exception of CAS\_750\_K\_1:1, which is a mixed Type I and IV isotherm with a hysteresis loop de-

**Table 2.** Selected textural parameters and yields of the chemically (KOH) activated carbon mono-**Table 2.** Selected textural parameters and yields of the chemically (KOH) activated carbon monoliths.


<sup>a</sup> Calculated at p/p<sup>0</sup> = 0.99; <sup>b</sup> calculated with the DR method; <sup>c</sup> <sup>V</sup>meso = Vtotal <sup>−</sup> <sup>V</sup>micro; d expressed as (massmonolith/masswhey) × 100.

The effect of the monolith/CH ratio on the KOH activation was as expected. For a given temperature (750 ◦C), increasing the KOH content increases the total pore volume of the monoliths from 0.655 to 1.224 cm3/g. The microporosity is enhanced as one moves from CAS\_750\_K\_1:1 to CAS\_750\_K\_1:3, whereas the mesopore volume diminishes quickly (Table 2). A close look to the PSD of these materials (Figure 7b) shows how the initially wide mesopores (>5 nm) present in CAS\_750\_K\_1:1 disappear as the relative CH ratio increases (i.e., samples CAS\_X\_K\_1:2 and CAS\_X\_K\_1:3, with X being 750 or 800), with narrow mesopores with an average diameter of around 2.8 nm evolving in those samples. The origin of the wide mesopores in sample CAS\_750\_K\_1:1 should be related to the precursor porosity [45]. Since the carbonised whey CW\_750 shows no significant N<sup>2</sup> adsorption at cryogenic temperatures, the origin of the wide mesoporosity of CAS\_750\_K\_1:1 should be related to the washing carried out after the KOH activation to remove byproducts (mainly carbonates). To check this possibility, a CW\_750 monolith was thoroughly washed with HCl and water following the same methodology used for the KOH-activated monoliths. The N<sup>2</sup> adsorption isotherm and textural parameters of the CW\_750 \_D monolith are also included in Figure 7 and Table 2. It is clear that (i) a simple demineralisation process increases significantly the porosity (as determined by the N<sup>2</sup> adsorption) of the whey carbon monoliths, with mineral species in the CWs blocking the entrance to pores; and (ii) mesopores are a significant part of CW\_750 \_D porosity, with a PSD resembling that of CAS\_750\_K\_1:1 (see inset in Figure 7b).

As for the effect of the activation temperature [45], for a given monolith/CH ratio (1:2), results fulfilled the expectances only partially. Increasing the heat treatment temperature from 700 ◦C to 750 ◦C leads to an important increment in the porosity of the monoliths (Figure 7 and Table 2). However, a further rise to 800 ◦C has little effect in either the specific surface area or pore volumes.

The effectiveness of the KOH activation when compared to thermal activation is clearly demonstrated with materials reaching SBET values up to 2815 m2/g and a Vtotal = 1.224 cm3/g for the maximum chemical agent ratio tested (1:3). Moreover, the yields of the activated monoliths per specific surface area are much higher in the case of the KOH-activated samples (Tables 1 and 2). In spite of this latter observation, the harshness of the alkali activation caused at least the presence of cracks that where evident after ocular inspection of the monoliths (Figure 5). In some cases, CAS\_750\_K\_1:3 and CAS\_800\_K\_1:2, the breakdown of the monoliths in several pieces occurred, thus limiting strongly the feasibility of this type of activation to obtain customised 3D microporous carbons.
