3.3.2. H3PO<sup>4</sup> Activation

As in the case of the KOH activation, we explored the effect of the CH relative proportion and the activation temperature variables on the outcome. Figure 8 shows the N<sup>2</sup> adsorption isotherms (−196 ◦C) of the H3PO<sup>4</sup> activated samples, with Table 3 collecting the principal textural parameters calculated. All phosphoric activated monoliths isotherms are a combination of Type I and Type IV for microporous materials with a certain amount of mesopores and hysteresis loop type H4 for slit-shaped pores. The isotherms have open "knees" in the micropore relative pressure range, which indicates a lower contribution of microporosity to the total pore volume. This contrasts with the KOH-activated samples showing isotherms with very close "knees". Thus, comparing the samples with similar Vtotal values (i.e., CAS\_750\_K\_1:1 and CAS\_800\_P\_1:2, Tables 2 and 3, respectively), the KOH-activated samples exhibit significantly higher SBET and Vmicro than the H3PO4 activated samples. As a consequence, the contribution of mesoporosity to the Vtotal is the highest for this type of activation. The N<sup>2</sup> PSDs of all the phosphoric-activated samples (Figure 8b) point out that a relatively sharp, unimodal distribution in the mesoporosity range is common to all monoliths, with maxima located at ca. 3 nm (i.e., narrow mesopores). *C* **2023**, *9*, x FOR PEER REVIEW 11 of 18 to the Vtotal is the highest for this type of activation. The N2 PSDs of all the phosphoric-activated samples (Figure 8b) point out that a relatively sharp, unimodal distribution in the mesoporosity range is common to all monoliths, with maxima located at ca. 3 nm (i.e., narrow mesopores).

**Figure 8.** (**a**) N2 adsorption isotherms (−196 °C); and (**b**) DFT pore size distributions of the chemically (H3PO4) activated monoliths derived from whey. **Figure 8.** (**a**) N<sup>2</sup> adsorption isotherms (−196 ◦C); and (**b**) DFT pore size distributions of the chemically (H3PO<sup>4</sup> ) activated monoliths derived from whey.

**Table 3.** Selected textural parameters and yields of the chemically (H3PO4) activated carbon mono-

CAS\_700\_P\_1:2 891 0.511 0.303 0.208 1.83 31.8 CAS\_800\_P\_1:2 1354 0.751 0.438 0.313 2.00 23.7 CAS\_800\_P\_1:3 886 0.539 0.297 0.242 1.93 26.4 CAS\_850\_P\_1:2 1062 0.578 0.372 0.206 1.93 22.1

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

The possibility of using different thermally stabilised whey:H3PO4 ratios was very limited and restricted in practice to a 1:2 ratio. The thermally stabilised whey: H3PO4 ratios of 1:1 produced little to no pore development, whereas the use of higher ratios (1:3) led to the overactivation of the monoliths. This latter observation is exemplified with the samples CAS\_800\_P\_1:2 and CAS\_800\_P\_1:3 in Figure 8 and Table 3. Increasing the relative proportion of the CH brings about a monolith with lower textural parameters. It has been reported that, when using high phosphoric acid impregnation ratios at relatively high (800 °C) activation temperatures, the progressive formation of polyphosphates and C-O-P bonds could block the pore entrances and lead to some loss of porosity, since these phosphate-like species are not easily removed in the washing process [46]. This would also justify the higher yield of CAS\_800\_P\_1:3 when compared to that of CAS\_800\_P\_1:2

Regarding the effect of the activation temperature, although the activation with H3PO4 works better at lower temperatures for lignocellulosic precursors [47,48], in this case, results show that higher temperatures (800 °C) are necessary to obtain an optimal porous development (Figure 8), with a maximum SBET for sample CAS\_800\_P\_1:2 of 1354

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

**ρHe (g/cm3)** 

Vmeso = Vtotal − Vmicro; d expressed as

**Yield d (%)** 

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

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

**(m2/g)** 

**Sample SBET**

(massmonolith/masswhey) × 100.

liths.

(Table 3).


**Table 3.** Selected textural parameters and yields of the chemically (H3PO<sup>4</sup> ) 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 possibility of using different thermally stabilised whey:H3PO<sup>4</sup> ratios was very limited and restricted in practice to a 1:2 ratio. The thermally stabilised whey: H3PO<sup>4</sup> ratios of 1:1 produced little to no pore development, whereas the use of higher ratios (1:3) led to the overactivation of the monoliths. This latter observation is exemplified with the samples CAS\_800\_P\_1:2 and CAS\_800\_P\_1:3 in Figure 8 and Table 3. Increasing the relative proportion of the CH brings about a monolith with lower textural parameters. It has been reported that, when using high phosphoric acid impregnation ratios at relatively high (800 ◦C) activation temperatures, the progressive formation of polyphosphates and C-O-P bonds could block the pore entrances and lead to some loss of porosity, since these phosphate-like species are not easily removed in the washing process [46]. This would also justify the higher yield of CAS\_800\_P\_1:3 when compared to that of CAS\_800\_P\_1:2 (Table 3).

Regarding the effect of the activation temperature, although the activation with H3PO<sup>4</sup> works better at lower temperatures for lignocellulosic precursors [47,48], in this case, results show that higher temperatures (800 ◦C) are necessary to obtain an optimal porous development (Figure 8), with a maximum SBET for sample CAS\_800\_P\_1:2 of 1354 m2/g (Table 3). This could be attributed to the different mechanisms involved in the activation with phosphoric acid. In the activation at low temperatures (500 ◦C) of the thermally stabilised (250 ◦C, in this case) whey monoliths, the porosity is mainly produced by the elimination of volatile matter (Figure 3) catalysed by the dehydration of the phosphoric acid, which promotes bond cleavage reactions and crosslinking formation via processes such as cyclisation and condensation [47,48]. Activation at higher temperatures (>500 ◦C) provokes, in some cases, the collapse of the porous network. However, when the activation temperature rises to 700 ◦C, the structure of the whey carbon is strong enough to maintain the microporosity developed at earlier stages (hence the similar pore development of samples CAS\_500\_P\_1:2 and CAS\_700\_P\_1:2), with a slight densification taking place as shown by the significant increase in the real density (ρHe) of the CAS\_700\_P\_1:2 sample.

As just mentioned, the most important pore development occurs at 800 ◦C, sample CAS\_800\_P\_1:2. The activation mechanism of phosphoric acid at such relatively high temperatures is different to that already discussed. At this temperature, H3PO<sup>4</sup> dehydrates and transforms into P4O<sup>10</sup> [40]. This latter moiety behaves in the absence of water as an oxidant reacting with the carbon matrix as follows [48]:

$$\rm P\_4O\_{10} + \rm 2C \rightarrow P\_4O\_6 + \rm 2CO\_2 \tag{1}$$

The formation of volatile phosphorous-containing compounds as a result of the present phosphate reduction and the introduction of phosphorus species into the carbon structure through C-O-P bonds led to the formation of new pores and the widening of the existent pores. At the same time, the reaction between the evolved CO<sup>2</sup> and the carbon skeleton results in an additional increase in porosity. Depending on the precursors, some authors report the presence of PH<sup>4</sup> and elemental P<sup>4</sup> as products of the high-temperature phosphoric activation of carbons [48–50]. These phosphorous species are easily removed in the washing process thus leaving pores available. Finally, further rising the activation temperature up to 850 ◦C (sample CAS\_850\_P\_1:2) backtracks in terms of the pore development, thus suggesting that some of the micro- and narrow mesoporosity start to collapse.

Although the porous development of the H3PO4-activated samples is poorer compared with the KOH-activated samples, the yields are considerably higher, higher even than the thermally activated ones. This would make this activation very interesting, although we could not obtain samples free of cracks (Figure 5) under the activation conditions tested.

### *3.4. Further Characterisation of Selected Samples*

Five monolithic samples combining high specific surface area and (in principle) good mechanical integrity, namely CW\_850, CW\_750\_D, TAW\_850\_1.5, CAS\_750\_K\_1:2, and CAS\_800\_P\_1:2 were further characterised using different techniques.

To complete the characterisation of the porosity of the monoliths, the selected samples were analysed by Hg intrusion. Figure 9 shows the PSD obtained by Hg intrusion with the corresponding main textural parameters collected in Table 4. All monoliths show an important macroporous development with the maxima of the Hg PSD centred at ca. 27 µm, except in the case of TAW\_850\_1.5, for which the Hg PSD is slightly shifted to smaller pore sizes values (Table 4). The percentage of porosity of the activated samples is >70%, with sample CAS\_750\_K\_1:2 showing an outstanding 82% porosity value. The widening of the PSD of CAS\_750\_K\_1:2 in the lower pore sizes region is also evident. This might suggest that the intrusion of Hg is damaging the sample. The same observation also stands when analysing the PSD of the monoliths in the mesoporous region. The relatively high Vmeso (as determined with Hg intrusion) of CAS\_750\_K\_1:2 was totally unexpected in the light of its N<sup>2</sup> adsorption isotherm (Figure 7). Actually, the only monoliths that show significant Vmeso values in the Hg intrusion characterisation are the activated ones (Table 4), thus pointing out that the monoliths might not be bearing such massive pressures. *C* **2023**, *9*, x FOR PEER REVIEW 14 of 18

**Figure 9.** Pore size distributions (PSD) of selected monoliths as measured by Hg intrusion. **Figure 9.** Pore size distributions (PSD) of selected monoliths as measured by Hg intrusion.

*3.5. Mechanical Properties of Selected Monoliths*  Mechanical properties of the selected samples were evaluated by compressive tests. Figure 10a shows typical strain/stress curves of the tested materials. Two of the monoliths, the chemically activated CAS\_750\_K\_1:2 and CAS\_800\_P\_1:2, could not bear any load during the compressive test, i.e., the structures shattered as soon as the load started to build up. The crevices on the surface of the chemically activated monoliths acted as fracture initiators that propagated quickly and destroyed the structures. Conversely, the compressive test curves of the CW\_850 and CW\_850\_D monoliths increased monoton-The elemental analyses, moisture and ash contents of the selected monoliths are collected in Table 5. Both CW\_850 and TAW\_850\_1.5 samples have the highest ash contents. The demineralisation of the carbonised samples and the cleaning of the by-products after the chemically activated monoliths remove most of the ashes leaves only <4% leftover. Those two samples also contain the highest moisture values and oxygen contents, and the lowest carbon contents, all of which being related to their high ash content. It is also remarkable the high nitrogen content that remains in the carbon structure of the activated monoliths (≥3%), regardless of the activation procedure. Specifically, chemical activation

ically within the elastic linear zone of the curve until the experiment reached the point of

monolith TA\_850\_1.2. In this case, the curve progresses with some breaks or jumps, which could indicate small fractures or imperfections in the monolith. After these discontinuities in the curve, the stress recovers almost immediately and the monoliths is capable of bearing loads up to higher strain values, when compared to CW\_850 and CW\_850\_D. The imperfections/fractures change the monolith failure mode to non-brittle, and if they exacerbate, as in the case of the chemically activated monoliths and the thermally activated monoliths under more severe conditions, thus leading to surface crevices

CW\_850 CW\_850\_D TAW\_850\_1.5

60 Su (MPa)

E (MPa)

E (MPa)

(Figure 5), the monolith cannot bear a substantial load.

(**a**) (**b**)

0

10

20

30

Su (MPa)

40

50

0.00 0.01 0.02 0.03 0.04 0.05 0.06

Strain

0

10

20

30

Stress (MPa)

40

50

60

 CW\_850 CW\_850\_D TAW\_850\_1.5

had little effect on the final nitrogen content of the monoliths. This is especially relevant for a broad spectrum of electrochemical applications with special interest in natural and induced nitrogen carbon materials [51–53]. This is very interesting since all this nitrogen content is naturally present in the whey powders, i.e., there is no need for an extra nitrogen source for N-doping. Furthermore, the N content of these materials surpasses even those reported for N-doped mesoporous carbons derived from tannins, as well as the microporous carbons derived from various natural biowaste fibres that have been modified with different nitrogen source species, such as urea or melamine [54,55]. Finally, monolithic materials with very different surface chemistry in terms of acidity/basicity, as determined by the pHPZC, were obtained.


**Table 4.** Textural parameters of selected monoliths obtained from Hg intrusion data.

<sup>a</sup> Volume of intruded Hg; <sup>b</sup> volume of Hg intruded in pores with sizes > 50 nm; <sup>c</sup> <sup>V</sup>meso = VHg <sup>−</sup> <sup>V</sup>macro; <sup>d</sup> pore size of the PSD maximum; <sup>e</sup> open porosity, s = [1 <sup>−</sup> (ρHg/ρHe)] <sup>×</sup> 100.


**Table 5.** Elemental analysis, moisture and ash contents, and pHPZC of selected monoliths.

<sup>a</sup> Elemental analysis, dry basis; <sup>b</sup> proximate analysis, wet basis.

#### *3.5. Mechanical Properties of Selected Monoliths*

Mechanical properties of the selected samples were evaluated by compressive tests. Figure 10a shows typical strain/stress curves of the tested materials. Two of the monoliths, the chemically activated CAS\_750\_K\_1:2 and CAS\_800\_P\_1:2, could not bear any load during the compressive test, i.e., the structures shattered as soon as the load started to build up. The crevices on the surface of the chemically activated monoliths acted as fracture initiators that propagated quickly and destroyed the structures. Conversely, the compressive test curves of the CW\_850 and CW\_850\_D monoliths increased monotonically within the elastic linear zone of the curve until the experiment reached the point of failure. This behaviour is characteristic of ceramic materials presenting brittle failure. Finally, a clear difference is observed in the stress/strain curve of the thermally activated monolith TA\_850\_1.2. In this case, the curve progresses with some breaks or jumps, which could indicate small fractures or imperfections in the monolith. After these discontinuities in the curve, the stress recovers almost immediately and the monoliths is capable of bearing loads up to higher strain values, when compared to CW\_850 and CW\_850\_D. The imperfections/fractures change the monolith failure mode to non-brittle, and if they exacerbate, as in the case of the chemically activated monoliths and the thermally activated monoliths under more severe conditions, thus leading to surface crevices (Figure 5), the monolith cannot bear a substantial load.

(Figure 5), the monolith cannot bear a substantial load.

**Figure 10.** (**a**) Examples of stress/strain curves; and (**b**) strength (Su) and modulus (E) of selected carbon monoliths. carbon monoliths.

Figure 10b shows the average values (and standard deviations) of the compressive

**Figure 9.** Pore size distributions (PSD) of selected monoliths as measured by Hg intrusion.

Mechanical properties of the selected samples were evaluated by compressive tests. Figure 10a shows typical strain/stress curves of the tested materials. Two of the monoliths, the chemically activated CAS\_750\_K\_1:2 and CAS\_800\_P\_1:2, could not bear any load during the compressive test, i.e., the structures shattered as soon as the load started to build up. The crevices on the surface of the chemically activated monoliths acted as fracture initiators that propagated quickly and destroyed the structures. Conversely, the compressive test curves of the CW\_850 and CW\_850\_D monoliths increased monotonically within the elastic linear zone of the curve until the experiment reached the point of failure. This behaviour is characteristic of ceramic materials presenting brittle failure. Finally, a clear difference is observed in the stress/strain curve of the thermally activated monolith TA\_850\_1.2. In this case, the curve progresses with some breaks or jumps, which could indicate small fractures or imperfections in the monolith. After these discontinuities in the curve, the stress recovers almost immediately and the monoliths is capable of bearing loads up to higher strain values, when compared to CW\_850 and CW\_850\_D. The imperfections/fractures change the monolith failure mode to non-brittle, and if they exacerbate, as in the case of the chemically activated monoliths and the thermally activated monoliths under more severe conditions, thus leading to surface crevices

*3.5. Mechanical Properties of Selected Monoliths* 

0

1

2

3

Log Differential Intrusion (cm3/g)

4

5

 CW\_850 CW\_750\_D TAW\_850\_1.5 CAS\_750\_K\_1:2 CAS\_800\_P\_1:2

10 100 1000 10000 100000

Pore Size Diameter (nm)

Figure 10b shows the average values (and standard deviations) of the compressive strength (Su) and modulus (E) of CW\_850, CW\_750\_D, and TAW\_850\_1.5. The monolith elastic modulus goes down with porosity, from 1.5 GPa to 350 MPa for CW\_850\_D and TAW\_850\_1.5, respectively. This latter value is, however, much higher than the elastic modulus reported for hierarchical carbon monoliths made from glucose (1 MPa) [56]. In the case of CW\_850, the S<sup>u</sup> value is over 40 MPa, which is superior to other porous carbons and even ceramic materials with similar porosities. As the porosity of the monoliths increases, the compressive strength decreases (Figure 11). In this way, CW\_850\_D presents values around 20 MPa, while for TAW\_850\_1.5 the value of S<sup>u</sup> drops to ca. 10 MPa. However, all these values are still very remarkable and considerably higher than those of other activated carbons found in the literature (Figure 11, red symbols) [19,20,57–62], specifically those made from coal (red triangles) [19], or those that are resin-based [57]. The red star of Figure 11 corresponds to a ceramic (cordierite)/carbon monolith [62]. strength (Su) and modulus (E) of CW\_850, CW\_750\_D, and TAW\_850\_1.5. The monolith elastic modulus goes down with porosity, from 1.5 GPa to 350 MPa for CW\_850\_D and TAW\_850\_1.5, respectively. This latter value is, however, much higher than the elastic modulus reported for hierarchical carbon monoliths made from glucose (1 MPa) [56]. In the case of CW\_850, the Su value is over 40 MPa, which is superior to other porous carbons and even ceramic materials with similar porosities. As the porosity of the monoliths increases, the compressive strength decreases (Figure 11). In this way, CW\_850\_D presents values around 20 MPa, while for TAW\_850\_1.5 the value of Su drops to ca. 10 MPa. However, all these values are still very remarkable and considerably higher than those of other activated carbons found in the literature (Figure 11, red symbols) [19,20,57–62], specifically those made from coal (red triangles) [19], or those that are resin-based [57]. The red star of Figure 11 corresponds to a ceramic (cordierite)/carbon monolith [62].

**Figure 11.** Correlation between Su and SBET of the selected carbon monoliths made from whey (black squares). The red symbols are results of other activated carbon monoliths found in the literature: (Liu, 2006) [19]; (Rangel-Sequeda, 2022) [20]; (Du, 2020) [57]; (Li, 2024) [58]; (Ibeh, 2019) [59]; (Guo, 2019) [60]; (Tang, 2020) [61]; (Gadkaree, 1998) [62]. **Figure 11.** Correlation between S<sup>u</sup> and SBET of the selected carbon monoliths made from whey (black squares). The red symbols are results of other activated carbon monoliths found in the literature: (Liu, 2006) [19]; (Rangel-Sequeda, 2022) [20]; (Du, 2020) [57]; (Li, 2024) [58]; (Ibeh, 2019) [59]; (Guo, 2019) [60]; (Tang, 2020) [61]; (Gadkaree, 1998) [62].

chined for specific applications. They are easily produced by pseudo-sintering dehydrated whey powders at high temperatures, which promote self-reaction between the precursor components (specifically lactose and whey proteins) with no need for external binders, overpressure, or templates. A mechanically strong porous carbon 3D structure with a hierarchical macro/meso/microporosity can be obtained simply by demineralising the whey carbon monoliths. The monolithic shape is preserved upon thermal activation with CO2 and chemical activation with KOH or H3PO4, resulting in high porosities and SBET up to 2400 m2/g, although chemical activation affects critically the mechanical properties of the 3D structures. The final carbonised, demineralised, and thermally activated monoliths have exceptional mechanical properties and an interconnected porous struc-

ture. In addition, they have a notable natural nitrogen content of up to 3 wt.%.

**4. Conclusions** 

#### **4. Conclusions**

A simple and novel method is presented for synthesising highly porous 3D carbon structures. These carbon monoliths can be easily moulded to specific shapes and machined for specific applications. They are easily produced by pseudo-sintering dehydrated whey powders at high temperatures, which promote self-reaction between the precursor components (specifically lactose and whey proteins) with no need for external binders, overpressure, or templates. A mechanically strong porous carbon 3D structure with a hierarchical macro/meso/microporosity can be obtained simply by demineralising the whey carbon monoliths. The monolithic shape is preserved upon thermal activation with CO<sup>2</sup> and chemical activation with KOH or H3PO4, resulting in high porosities and SBET up to 2400 m2/g, although chemical activation affects critically the mechanical properties of the 3D structures. The final carbonised, demineralised, and thermally activated monoliths have exceptional mechanical properties and an interconnected porous structure. In addition, they have a notable natural nitrogen content of up to 3 wt.%.

The possibility of using whey as a precursor of activated carbon monoliths constitutes a modest alternative for whey valorisation, in addition to lessening their dependence on fossil-fuel derived precursors. These two outcomes are significant in a circular economy model of materials/goods production and consumption.

**Author Contributions:** Conceptualization, L.A.R.-M., J.A.M. and M.A.M.-M.; methodology, all authors; investigation, R.L.-U. and L.A.R.-M.; writing—original draft preparation, L.A.R.-M.; writing—review and editing, all authors; funding acquisition, J.A.M. and M.A.M.-M. All authors have read and agreed to the published version of the manuscript.

**Funding:** This research was funded by the Spanish Ministerio de Ciencia e Innovación (MCIN/AEI/ 10.13039/501100011033) (Project PID2020-115334GB-I00) and Principado de Asturias (FICYT)—European Union (FEDER) (Project PCTI-Asturias IDI/2021/000015). L.A.R.-M. thanks CONACYT, Mexico, for a postdoctoral grant (CVU No. 330625, 2022).

**Data Availability Statement:** The data presented in this study are available on request from the corresponding author.

**Acknowledgments:** J. Angel Menéndez and Miguel Montes are members of the CSIC Interdisciplinary Thematic Platform (PTI+) for Sustainable Plastics towards a Circular Economy (PTI-SusPlast+).

**Conflicts of Interest:** The authors declare no conflict of interest. The funders had no role in the design of the study; in the collection, analyses, or interpretation of the data; in the writing of the manuscript; or in the decision to publish the results.

#### **References**


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