vibrating lattice ($\rho_{lattice}$), which has a strong temperature dependence. However, microstructural features, especially in metallic alloys, can increase the electrical resistivity. Matthiessen's rule expressed in Equation 12 considers the contribution of two components: (i) lattice resistivity ($\rho_{lattice}$) and (ii) residual resistivity ($\rho_{residual}$). The latter is associated with local distortion of the lattice, solute atoms, grain boundaries and crystal defects. Their contribution to the total resistivity can be added and designated as B, which is weakly dependent on temperature. A is a constant independent of temperature [852–854].

$$ \rho_E = \rho_{lattice} + \rho_{residual} \approx A \cdot T + B \quad (12) $$

Processing copper and aluminum alloys by SPD increases their strength and decreases the electrical conductivity due to the high density of grain boundaries and crystal defects, which hinder dislocation motion but promote electron scattering mechanisms [180, 236, 237].

<figure><img src="image_111.png" /></figure>

<figure><img src="image_112.png" /><figcaption>Fig. 40. Electrical conductivity with respect to strength in (a) copper and (b) aluminum alloys [834].</figcaption></figure>

855–862]. Traditionally, high strength and highly conductive Cu-based alloys are expected to have ultimate tensile strength (UTS) of 300–800 MPa and electrical conductivity of 50–95 IACS% [856]. Meanwhile, Al-based alloys should achieve at least a UTS of 300 MPa and 52 IACS% [857]. Such ranges of properties are indicated in Figs. 40a and 40b as hatched areas. Copper alloys can be strengthened with restored electrical conductivity through several processing routes. A summary of such results is shown in Fig. 40a.

* Grain refinement by HPT of high-purity copper increases its strength to 440 MPa [859] with electrical conductivity of ~87 IACS% at the steady state [860]. During microstructure evolution high density of dislocations are produced, but the microstructure undergoes dynamic recrystallization with a grain size of ~ 0.2 µm and lower density of dislocations at the steady state.

* Age hardenable Cu-Cr-X alloys processed by SPD with further annealing allow the precipitation of particles rich in chromium. Meanwhile, SPD at high temperatures promotes dynamic strain aging and particle precipitation. The low solute content in the copper matrix reduces the electron scattering and optimizes the electrical conductivity [254, 856, 861–863]. Islamgaliev et al. [862] processed a Cu-Cr-Ag alloy by HPT at room temperature followed by aging at 723 K for 1 h and achieved UTS of 840 MPa and conductivity of 85 IACS%, as shown in Fig. 40a.

* Fine grains and nanoscale twins have been shown to minimize the scattering of conduction electrons and still serve as barriers to dislocation motion [864–868]. Fig. 41a shows a bright-field image of high-purity copper with nanoscale twins produced by pulsed electrodeposition (PED), which shows yield stress of ~900 MPa and conductivity of 96.9 IACS%. This result is also shown in Fig. 40a [864].

The mechanical and electrical properties of aluminum alloys have a close relationship with the nature and fraction of alloying elements [834, 855, 870]. UFG Al-based alloys subjected to post-processing heat treatment after SPD or SPD at elevated temperatures have been reported to produce decomposition of supersaturated solid solutions and precipitation of second phase particles of Mg₂Si [870–876], Al₃Fe, Al₆Fe [869, 877–880], Al₁₁RE₃ (RE: rare-earth elements) [881, 882] and Al₂Zr [332, 883–886]. A summary of these results is shown in Fig. 40b. For instance, processing AA6201 by HPT at elevated temperatures (403 K, 453 K and 503 K) allowed the precipitation of nano-sized Mg₂Si due to dynamic strain aging and increased electrical conductivity with processing temperature: 55.6 IACS%, 58.4 IACS% and 59.0 IACS% respectively [870]. Consequently, strength decreases to 412, 365 and 275 MPa by annealing, respectively, as shown in Fig. 40b. On the other hand, UFG Al-2Fe (wt%) alloys with nanosized Al₃Fe and Al₆Fe particles, as indicated by arrows in Figs. 41b and 41c were produced by HPT for 20 and 75 revolutions, respectively. Both were peak-aged at 473 K and attained UTS of 600 and 650 MPa with 55 and 54 IACS% respectively [869, 877]. This result is also shown in Fig. 40b.

In summary, the thermal dependence of electrical resistivity, grain boundaries and grain size allows the balance of mechanical and electrical properties by thermomechanical routes, such as SPD complemented by static or dynamic aging [9, 373, 887]. On the other hand, scaling up sample dimensions and optimization of continuous SPD methods can open a new path for the industrial application of high-performance UFG copper and aluminum conductors [154, 164, 179, 241].

## 6.2. Magnetic properties

Magnetic materials are key materials needed, for the electrification of mobility and the generation and conversion of electrical energy [888]. In recent years, interest in the magnetic properties of SPD-processed materials increased as microstructure strongly influences
<figure><img src="image_113.png" /><figcaption>Fig. 41. (a) TEM bright-field image of as-deposited copper with diffraction pattern (as inset) [864]. (b,c) TEM bright/dark-field images of Al-2Fe (wt%) alloy after HPT processing for (b) 20 and (c) 75 revolutions at peak-aged condition [869].</figcaption></figure>

<figure><img src="image_114.png" /><figcaption>Fig. 42. (a) Schematic representation of coercivity ($H_c$) as a function of grain size ($D$) for soft (upper graph) and hard magnetic (lower graph) materials (different logarithmic scales for $H_c$ of soft and hard magnetic are used) [888–893]. (b) Powder blends for the HPT process, (c) schematic representation of powder-based HPT processes and (d) schematics of bulk materials produced by HPT from powder blends with enhanced microstructural features [841].</figcaption></figure>

the magnetic properties and HPT can be used to tune it [841]. For example, the coercivity $H_c$ of soft and hard magnetic materials correlates with the grain size (Fig. 42a) [888–893] and HPT leads to strong grain refinement [23, 38, 894]. Other advantages include the potential to form nanostructured composites with a high density of interfaces or supersaturated solid solutions, while at the same time enabling a shear

texture formation through the deformation processes [895, 896]. Powder-based HPT processing (Fig. 42b-d) offers even more interesting and unique opportunities for the fabrication of hard or soft magnetic materials [841]. Using elemental powders or intermetallic phases in the form of powders, a free choice of combinations is possible, without the limitations given by the phase diagrams or, better, by phase formation
<figure><img src="image_115.png" /><figcaption>**Fig. 43.** Different types of microstructures which can be generated by HPT applied to powder/powder blends and their corresponding schematic magnetic hysteresis curves [841].</figcaption></figure>

from a melt as in metallurgical processing routes (Fig. 42b). As mentioned earlier, HPT processing of powders can produce bulk nano-structured materials (e.g. nanocomposites), while at the same time processes like mechanical alloying, phase transformations and super-saturation can be used to enhance the magnetic properties [897, 898]. The free choice of starting powders is particularly interesting because thermodynamically metastable chemical compositions and various composite structures can be achieved. Fig. 43 summarizes examples of microstructures of different magnetic materials or nanocomposites that can be obtained by HPT processing of powders [841]. Here some recent findings are reviewed.

Smaller grain sizes compared to single-phase materials can be achieved in HPT-deformed Cu-Co [897–899] and Fe-Cu alloys [900] with supersaturated phases, in which semi-hard, tunable magnetic properties are obtained. Additionally, annealing treatments are used to modify magnetic properties [899]. To achieve soft magnetic properties, ternary Fe-Co-Cu alloys were studied [901]. The addition of iron decreases the magnetocrystalline anisotropy and increases the saturation magnetization [902]. Giant magneto-resistance (GMR) is observed in materials that consist of nanometer-sized ferromagnetic particles dispersed in non-magnetic metallic matrixes (termed ‘granular’) [903, 904]. In HPT-deformed binary and ternary powder mixtures consisting of different compositions of elemental copper, cobalt, nickel and iron [905–907], similar GMR values as in materials processed with different methods are found. HPT processing of powders is further used to obtain hard magnetic materials. In recent studies, Nd-Fe-B-based bulk materials were processed by HPT from amorphous precursors, in which the deformation-induced texture further led to anisotropic magnetic properties and a maximum energy product of 22.8 MGOe [908, 909].

For nucleation-controlled magnets, the four classical magnetic interaction mechanisms with the different microstructures (type I-IV) can be implemented via powder-based HPT (Fig. 43). Powder blends of SmCo₅ and magnetically different binder phases, i.e. copper (diamagnetic) [910], iron (ferromagnetic) [911], nickel (ferro), cobalt (ferro), zinc (para), tin (para) and chromium (antiferromagnetic) [841], were deformed by HPT. A type I magnetic material (aligned/textured single-domain ferromagnetic particles are separated by a thin non-magnetic layer leading to magnetic decoupling of ferromagnetic particles) was obtained using copper as a diamagnetic phase [910]. Type II magnets (exchange-coupled ferromagnetic particles) were fabricated as well by deforming SmCo-based intermetallic powders by HPT [841]. An exchange-coupled type III system (a combination of soft and hard magnetic phases) was synthesized using iron in addition to SmCo₅ [911]. A type IV ferro-antiferromagnetic composite is processed by deforming SmCo₅ and chromium [841]. The highest saturation

magnetization was obtained for the Fe-containing composition (type III) and the lowest for the diamagnetic and antiferromagnetic binder phases (type I and IV). According to Fig. 42a it has been shown that a refinement of the microstructure and grain size of the hard ferromagnetic phase increases the coercivity, whereas the largest coercivity is achieved with copper and chromium as a binder phase [841]. The highest coercivity is obtained in the Cu-SmCo₅ system (type I) due to the magnetic decoupling of the SmCo₅ particles by copper [910]. Successful processing of type IV magnetic materials by HPT was also reported in the Fe-NiO [912] and Co-NiO [913] systems.

Phase transformations by HPT can also be employed to synthesize different hard magnetic phases such as L1₀-FeNi [340], Mn-Al-C [914] and MnBi [915]. Soft magnetic properties in HPT-processed bulk single-phase materials (cobalt [509, 916], nickel [917–919], iron [919–923], Fe-Si [917, 919, 921], FeCo [917, 921], FeNi [921]) have not yet been achieved due to large grain sizes and high density of defects. In HPT-processed hard magnetic materials, grain refinement generally leads to coercivity enhancement [924–928], partial amorphization [924, 929–932] or phase decomposition [933]. By deformation of amorphous magnetic materials [934–939], type III magnetic materials can be obtained [934].

In summary, SPD processing and particularly powder-based HPT offers strong advantages in the design of magnetic microstructures based on the possibility to freely combine hard, soft and non-magnetic phases for forming a textured nanocomposite. SPD processing opens thereby new pathways to future magnetic materials with unique microstructures covering the whole range of soft-to-hard magnetic properties.

### 6.3. Optical bandgap narrowing and solar energy harvesting

Metal oxides, such as TiO₂ and ZnO, are efficient semiconductors for light harvesting and photocatalysis because they generate a large number of electrons and holes upon light irradiation. However, their efficiency is low under sunlight because their bandgap is large and is in the range of ultraviolet (UV) light and sunlight only contains about 2–3% of UV light. To enhance the efficiency of metal oxide across the solar spectrum, it is necessary to broaden their light absorption to the visible light region by reducing their bandgap. Various approaches, such as doping with anions like nitrogen [940–942], carbon [943], hydrogen [944], fluorine [945], chlorine [945], bromine [945] and iodine [945] and/or cations like tungsten [946, 947], ruthenium [948], tantalum [949], sodium [950], bismuth [951], lanthanum [952], molybdenum [953], zinc [954], calcium [954] and cobalt [955] have been employed to reduce the bandgap of metal oxide semiconductors. When metal oxide is doped with anions like nitrogen or carbon, N 2p states in nitrogen and
C 2p states in carbon couple with O 2p states of the metal oxide, resulting in the formation of acceptor levels above the valence band of the host metal oxide [956]. The formation of such acceptor levels reduces the bandgap of the metal oxide whose valence band is composed of O 2p orbitals. Although doping or co-doping with anions/cations is an effective method for reducing the bandgap of most metal oxides, it has been challenging in some aspects. Dopants often tend to precipitate and alter the chemical composition of the host material. Another limitation is the low solubility of dopants in the host material. Additionally, high temperatures are typically required to enhance the crystallinity of the host material and this can lead to the unintended expulsion of dopants.

An effective method for reducing the bandgap is the introduction of defects into the host material. Taking TiO₂ as an example, defects such as oxygen vacancies have been shown to reduce its bandgap. The major intrinsic defects in TiO₂ include oxygen vacancies, titanium interstitials and titanium vacancies [957–959]. Among these, oxygen vacancies are the most common type of intrinsic defect. Introducing plastic strain by HPT method is shown to be effective for the formation of the oxygen vacancies in many semiconductors such as TiO₂ [582, 639, 960], ZnO [604], Y₂O₃ [583], BiVO₄ [961], TiO₂/ZnO [605], CsTaO₃ [962], LiTaO₃ [962], Bi₂O₃ [963], ZrO₂ [666], GaN/ZnO [606], BaTiO₃ [840] and Al₂O₃ [964]. In this method, dopants are not used, while oxygen vacancies are effectively formed. The generation of oxygen vacancies arises from the crystal structure of most oxides including TiO₂. In TiO₂, bridge bonding oxygen atoms are two-fold coordinated and can be easily removed from the lattice by introducing plastic strain. Therefore, oxygen vacancies are known to be the most reactive sites in TiO₂. Generation of one oxygen vacancy leaves two valence electrons that occupy empty Ti 3d levels. As a consequence, energy levels below the conduction band are generated, leading to the bandgap reduction of the metal oxide semiconductor.

Another interesting feature of introducing the simultaneous plastic strain and high pressure is reducing the bandgap of materials by stabilizing the high-pressure phases at ambient pressure. These high-pressure phases have theoretically been calculated to have better light absorbance and narrower bandgaps that coincide with visible light. For example, columbite TiO₂-II with an orthorhombic structure is calculated to have a bandgap of 2.59 eV, but it is stable only at pressures higher than 2 GPa [639, 960]. The application of HPT for stabilizing the TiO₂-II phase showed that the high-pressure phase was stable for at least 6 months. The high-pressure rocksalt phase of ZnO [604] and the monoclinic phase of Y₂O₃ [583] are some successful examples of bandgap reduction through the formation of high-pressure phases. Fig. 44a shows the UV-Vis light absorbance spectra of TiO₂ samples before HPT and after HPT. The absorption edge of the starting TiO₂ powder is 400 nm, while it shifted to 470 nm, after HPT processing. The bandgap of samples was calculated by the Kubelka-Munk equation as shown in Fig. 44b, being 3.1 eV and 2.7 eV for samples before HPT and after HPT processing. Comparing the bandgaps of samples before HPT and after HPT processing, it is clear that the formation of a high-pressure phase is an effective approach for reducing the bandgap of oxides. A schematic illustration of bandgap reduction is shown in Fig. 45. Fig. 45a represents the bandgap of a pure metal oxide, whose bandgap is shown to be wide. Fig. 45b shows the cation-doped oxide, where donor states are responsible for the bandgap narrowing. Fig. 45c shows an anion-doped oxide, where donor states are responsible for bandgap narrowing. Bandgap narrowing through oxygen vacancies produced by HPT processing also falls into this group, with the difference that no dopants are present in this case. Fig. 45d represents the bandgap of a pure metal oxide with high-pressure phases. In this case, the high-pressure phase itself has a narrower bandgap without any donor/acceptor states. Table 5 is a summary of publications in the literature that studied the effect of the HPT method on the bandgap of various semiconductors and ceramics [323–327, 844, 965, 966].

Taken altogether, inducing plastic strain using the SPD is an effective approach for the formation of oxygen vacancies and/or high-pressure

<figure><img src="image_116.png" /></figure>

<figure><img src="image_117.png" /><figcaption>Fig. 44. (a) UV-Vis absorbance spectra and (b) bandgap calculation by the Kubelka-Munk theory for TiO₂ before and after HPT processing and after HPT followed by annealing [639].</figcaption></figure>

<figure><img src="image_118.png" /><figcaption>Fig. 45. Schematic illustration of bandgap reduction mechanisms: (a) pure metal oxide, (b) cation-doped oxide in which donor states are responsible for bandgap narrowing, (c) anion-doped oxide in which acceptor states are responsible for bandgap narrowing and (d) high-pressure phases with oxygen vacancies in which bother donor and acceptor states are responsible for bandgap narrowing.</figcaption></figure>

phases in semiconductors to enhance light harvesting via bandgap narrowing. This feature of SPD has opened a pathway for the development of photocatalysts, as will be discussed in the next chapters.

## 6.4. Photocatalytic CO₂ conversion

Among all scientific attempts to solve the global warming crisis, photocatalytic CO₂ conversion has been introduced as one of the cleanest methods to convert CO₂ to CO and useful organic components [967–969]. Photocatalytic CO₂ conversion reactions are performed on the surface of a semiconductor under solar irradiation where electrons and holes are separated and immigrated to the surface as shown in Fig. 46a [970]. A variety of methods such as doping with impurities [971–974], defect/surface engineering [975, 976] and the introduction of heterojunctions [977, 978] have been reported to enhance the
<table><caption>**Table 5**
Summary of publications related to the effect of SPD via the HPT method on bandgap of various semiconductors and ceramics.</caption><thead><tr><th>Material</th><th>Bandgap before HPT (eV)</th><th>Bandgap After HPT (eV)</th><th>Factors Affect Bandgap</th><th>Color Before HPT</th><th>Color After HPT</th><th>Application</th><th>Ref.</th></tr></thead><tbody><tr><td>TiO<sub>2</sub></td><td>3.1</td><td>2.4</td><td>High-pressure phase, Oxygen vacancy</td><td>White</td><td>Green</td><td>Water splitting H<sub>2</sub> production</td><td>[582, 639, 960]</td></tr><tr><td>ZnO</td><td>3.4</td><td>1.8</td><td>High-pressure phase, Oxygen vacancy</td><td>White</td><td>Yellow</td><td>Dye degradation</td><td>[604]</td></tr><tr><td>Y<sub>2</sub>O<sub>3</sub></td><td>5.82</td><td>5.69</td><td>Phase transformation</td><td>White</td><td>Pink</td><td>Photoluminescence</td><td>[583]</td></tr><tr><td>BiVO<sub>4</sub></td><td>2.4</td><td>2.1</td><td>Defect formation</td><td>Yellow</td><td>Orange and Rose</td><td>CO<sub>2</sub> conversion</td><td>[961]</td></tr><tr><td>TiO<sub>2</sub>-ZnO</td><td>3.2</td><td>1.6</td><td>High-pressure phase, Oxygen vacancy</td><td>White</td><td>Black</td><td>Water splitting H<sub>2</sub> production</td><td>[605]</td></tr><tr><td>CsTaO<sub>3</sub></td><td>4.6</td><td>3.3</td><td>Oxygen vacancy</td><td>-</td><td>-</td><td>Water splitting H<sub>2</sub> production</td><td>[962]</td></tr><tr><td>LiTaO<sub>3</sub></td><td>4.7</td><td>4.2</td><td>Oxygen vacancy</td><td>-</td><td>-</td><td>Water splitting H<sub>2</sub> production</td><td>[962]</td></tr><tr><td>Bi<sub>2</sub>O<sub>3</sub></td><td>2.5</td><td>2.3</td><td>Oxygen vacancy</td><td>Yellow</td><td>Black</td><td>Photocurrent</td><td>[963]</td></tr><tr><td>ZrO<sub>2</sub></td><td>5.1</td><td>4.0</td><td>High-pressure phase, Oxygen vacancy</td><td>White</td><td>Black</td><td>Water splitting H<sub>2</sub> production</td><td>[666]</td></tr><tr><td>GaN-ZnO</td><td>2.7</td><td>2.4</td><td>Nitrogen vacancy</td><td>Yellow</td><td>Magenta</td><td>Water splitting H<sub>2</sub> production</td><td>[606]</td></tr><tr><td>Al<sub>2</sub>O<sub>3</sub></td><td>5.7</td><td>2.5</td><td>Oxygen vacancy</td><td>White</td><td>Black</td><td>Dye degradation</td><td>[964]</td></tr><tr><td>TiHfZrNbTaO<sub>11</sub></td><td>0</td><td>2.9</td><td>High-entropy phase</td><td>Metallic</td><td>Orange</td><td>Water splitting H<sub>2</sub> production, CO<sub>2</sub> conversion</td><td>[323, 324]</td></tr><tr><td>TiZrNbTaWO<sub>12</sub></td><td>&gt;3.0</td><td>2.3</td><td>High-entropy phase</td><td>White</td><td>Orange</td><td>Water splitting O<sub>2</sub> production</td><td>[325]</td></tr><tr><td>TiZrHfNbTaO<sub>6</sub>N<sub>3</sub></td><td>2.9</td><td>1.6</td><td>High-entropy phase</td><td>Orange</td><td>Black</td><td>Water splitting H<sub>2</sub> production, CO<sub>2</sub> conversion</td><td>[326, 327]</td></tr></tbody></table>

<figure><img src="image_119.png" /><figcaption>**Fig. 46.** (a) Schematic of photocatalytic CO₂ conversion mechanism [981]. (b) Photocatalytic CO production rate of BiVO₄ versus time for powder and samples processed by HPT with 0.25, 1 and 4 turns [961]. (c) High-resolution TEM image of sample processed by HPT containing TiO₂-II high-pressure phase [985]. (d) Rate of CO generation for TiO₂ samples before HPT, after HPT and after HPT and annealing [985].</figcaption></figure>

efficiency of different photocatalysts, but doping with impurities is the most feasible one. Since this method suffers from the impurity-induced recombination of electrons and holes [974, 979, 980], it is necessary to introduce new strategies to improve the photocatalytic activity. SPD through the HPT method which was employed to improve the photocatalytic dye degradation [604, 956, 964] and water splitting [606, 639, 666], showed high potential to suppress electron-hole recombination and enhance the photocatalytic CO₂ conversion efficiency. SPD enhanced photocatalytic CO₂ conversion by the introduction of (i) oxygen vacancies and strain, (ii) high-pressure phases, (iii) defective high-entropy oxides and (iv) low-bandgap high-entropy oxynitrides, as

discussed below [981, 982].

* Simultaneous oxygen vacancy generation and strain engineering using the HPT method was employed to improve the properties and photocatalytic activity of BiVO₄ for CO₂ conversion [961], due to the high recombination rate and inappropriate conduction band position of this material [983]. Using various characterization tools, it was observed that lattice strain and oxygen vacancies increase in the BiVO₄ sample by increasing the HPT turns [961]. These changes resulted in systematically increasing the light absorbance, narrowing the bandgap and suppressing the recombination rates. The
conduction band was also aligned to an upper level of energy which could advance the reactions thermodynamically. The BiVO₄ sample processed by HPT showed an improved CO production rate under UV light and the activity increased by increasing the HPT turns due to increasing lattice strain and oxygen vacancy as shown in Fig. 46b [961]. A similar concept was recently employed by HPT processing to produce oxygen vacancy-rich black rookite as an active photocatalyst for CO₂ conversion [984].

* As mentioned in the earlier chapter, columbite TiO₂-II is the high-pressure phase of TiO₂ which was formed by SPD through the HPT method (Fig. 46c) and employed for photocatalytic CO₂ conversion [985]. The resultant material after HPT was annealed to decrease the oxygen vacancy concentration in bulk. The annealed sample containing the TiO₂-II phase showed higher light absorbance in the visible region and lower bandgap. Recombination of electrons and holes suppressed significantly for the annealed sample and it could successfully generate a photocurrent much higher than the initial anatase powder. All these improvements led to considerable photocatalytic activity of anneal sample containing TiO₂-II compared to anatase powder as shown in Fig. 46d [985].

* Since high-entropy ceramics have shown superior properties for various ranges of applications, especially in the catalysis field [986, 987], a defective high-entropy oxide was fabricated using HPT and consequent high-temperature oxidation and examined for photocatalytic CO₂ conversion [324]. The synthesized high-entropy oxide TiZrNbHfTaO₁₁ containing nanograin boundaries and defects such as oxygen vacancy and dislocations adsorbed the light in both UV and visible regions with the bandgap value of 3 eV. It showed lower recombination of electrons and holes compared to anatase TiO₂ and BiVO₄ as conventional photocatalysts and successfully generated

photocurrent. This material showed a high CO production rate similar to P25 TiO₂ as a benchmark catalyst and lower than anatase TiO₂ and BiVO₄. This study reported the first application of high-entropy materials for photocatalytic CO₂ conversion [324].

* By considering the high potential of the mentioned TiZrNbHfTaO₁₁ synthesized by HPT for CO₂ photoreduction and the promising low bandgap of metal oxynitrides [988], the TiZrNbHfTaO₆N₃ high-entropy oxynitride was fabricated by nitriding the TiZrNbHfTaO₁₁ [327]. The material showed an appropriate electronic band structure with an extremely narrow bandgap of 1.6 eV which is lower than P25 TiO₂ and the high-entropy oxide as shown in Fig. 47a [327]. The light absorbance of this high-entropy oxynitride was higher than P25 TiO₂, high-entropy oxide and all corresponding binary oxides (Fig. 47b) [326, 327]. The recombination rate of this material was significantly suppressed compared with P25 TiO₂ and the high-entropy oxide and it adsorbs the CO₂ molecules with both physisorption and chemisorption better than P25 TiO₂ and the high-entropy oxide (Fig. 47c) [327]. The result of these improvements caused by HPT processing was the effective enhancement of photocatalytic CO₂ conversion activity of TiZrNbHfTaO₆N₃ compared to P25 TiO₂ as a benchmark photocatalyst, the high-entropy oxide (Fig. 47d) and almost all reported photocatalysts in the literature [327].

In summary, SPD was introduced as an effective tool to improve the structural and optical properties and consequently activity of conventional photocatalysts for CO₂ conversion. This method also showed promising potential to synthesize the new high-entropy ceramics with superior photocatalytic efficiency comparable to benchmark photocatalysts for CO₂ conversion. The ability of HPT to develop these highly

<figure><img src="image_120.png" /><figcaption>Fig. 47. (a) Electronic band structure of TiZrNbHfTaO₆N₃ high-entropy oxynitride (HEON), P25 TiO₂ and TiZrNbHfTaO₁₁ high-entropy oxide (HEO) [327]. (b) UV-vis light absorbance spectra for TiZrNbHfTaO₆N₃ compared to TiZrNbHfTaO₁₁ and corresponding binary oxides [326]. (c) Diffuse reflectance infrared Fourier transform spectra for TiZrNbHfTaO₆N₃ compared to P25 and TiZrNbHfTaO₁₁ [327]. (d) Rate of CO production for TiZrNbHfTaO₆N₃ compared to P25 TiO₂ and TiZrNbHfTaO₁₁ [327].</figcaption></figure>
active catalysts for CO₂ photoreduction is attributed to the formation of defects such as oxygen vacancies which act as active sites to adsorb and activate the CO₂ molecules [981, 982]. For the future outlook, the SPD can significantly extend in the field of photocatalysis to develop new and effective catalysts to convert CO₂ to CO and also hydrogenate to other useful components such as CH₄. The HPT method is expected to play the most important role in this regard because the method applies to a wide range of hard-to-deform materials such as optical ceramics [15, 844], semiconductors [512, 989] and oxides [513, 989].

## 6.5. Photocatalytic hydrogen production

Clean production of hydrogen is a promising alternative to fossil fuel

utilization, but its production is still a challenge [990]. Photocatalysis, by splitting water in the presence of light and a photocatalyst [991, 992], produces hydrogen with minimum CO₂ emission. Fig. 48a shows a schematic representation of a photocatalytic process. The main limitation of the photocatalysis is the photocatalyst itself. The photocatalyst must have great light absorbance, narrow bandgap, good efficiency and high stability. Several popular photocatalysts such as TiO₂ [639, 993], ZnO [604, 994, 995], perovskites [323, 962] and oxynitrides [606, 996] show these characteristics; however, various studies demonstrate the necessity of the improvement of activity of these catalysts. SPD processing via HPT [604, 639, 962], ball milling [997–1000] and tubular channel angular pressing (TCAP) [1001] were demonstrated to be beneficial in improving photocatalytic activity. SPD can tune three

<figure><img src="image_121.png" /><figcaption>**Fig. 48.** (a) Schematic representation of the photocatalytic water splitting reaction for hydrogen production. (b) Enhanced photocatalytic activity of zirconia by oxygen vacancy introduction via HPT processing [666]. (c) Enhanced photocatalytic activity TiO₂ by stabilization of high-pressure columbite phase [1002]. (d) Density functional theory calculations, showing high surface activity of the (101) atomic plane of high-pressure columbite for H₂O to OH + H decomposition [1002]. (e) Photocatalytic hydrogen production for TiO₂-ZnO composite under UV light after HPT processing for 3 and 15 turns [605]. (f) Schematic representation of charge separation in TiO₂-ZnO interface boundary as heterojunctions [605].</figcaption></figure>
major characteristics of the photocatalysts: (i) defects introduction [606, 666], (ii) high-pressure phase transformation [604, 639, 1002] and (iii) heterojunctions [325, 605, 997, 998].

Due to the pressure-strain effect, vacancy formation is one of the most common defects when processing ceramics by SPD. A large concentration of oxygen vacancies was found in black ZrO₂ [666], TiO₂ [639], ZnO [604, 1003], Al₂O₃ [995], CsTaO₃ [962] and LiTaO₃ [962] and MgO [956] after being processed by SPD. The oxygen vacancies are beneficial for photocatalytic activity in all the cases [604, 666, 956, 962, 995]. The study performed on black ZrO₂ demonstrated that oxygen vacancy concentration can be increased by high-temperature processing [666]. In addition, these strain-induced vacancies improve the hydrogen generation in the sample as shown in Fig. 48b. The oxygen vacancies can enhance the photon absorption, narrow optical bandgap and function as active sites for photocatalytic activity [639, 666, 956, 1002]. Nitrogen vacancies were also observed in GaN-ZnO after HPT [606]. The study concluded that the nitrogen-vacancy complexes can reduce the bandgap by modifying the energy levels above the top of the valence band and below the conduction band [606]. Vacancies generally diminish the electron-hole recombination and narrow the bandgap of the material. The defect introduction by SPD processing not only can improve several properties but also contribute to the stabilization of high-pressure phases.

The transformation to a high-pressure phase can be highly beneficial for photocatalytic activity. For ZnO, the transformation to Rocksalt decreases the bandgap from 3.4 to 1.8 eV [604]. However, as mentioned earlier, the rocksalt phase is only stable under high pressure but could be stabilized in ambient pressure by HPT [604]. Similarly, the TiO₂ high-pressure columbite phase is unstable under ambient pressure; however, it can be stabilized by the shear strain introduced by HPT [1004]. The shear strain induces gradual modifications on the anatase crystal structure until it transforms to ~70% columbite [1004]. Columbite stabilization in combination with the anatase, rutile and intermediate sheared structures significantly improves the water-splitting efficiency (Fig. 48c) [639, 1002, 1004]. Columbite is demonstrated to be highly active for water splitting in the atomic plane (101). The DFT results for the reaction pathway observed in Fig. 48d showed low activation energy for the water-splitting reaction compared to anatase and rutile [1002].

Several composites have been also subjected to plastic deformation or SPD methods to improve their coupling: SnO₂/TiO₂ [997, 1005], p-CuO/n-ZnO [998], p-CuBi₂O₄/n-TiO₂ [999], g-C₃N₄/ZnO [1000], CuFe₂O₄/Bi₄Ti₃O₁₂ [1006] and TiO₂-ZnO [605, 1005]. As shown in Fig. 48e, TiO₂-ZnO after HPT processing for 3 turns showed greater photocatalytic activity compared to TiO₂ and ZnO. Ball milling is commonly utilized to couple the oxides into a composite [997, 1006], but it is also able to generate improvements in the photocatalytic activity of the composites. At an optimum time of milling, the activity is enhanced, but it reduces after increasing the time due to the agglomeration of the powders [998, 999] or reduction of crystallinity [605]. The reason for the enhancement of the photocatalytic activity, as in the HPT, is attributed to the presence of heterojunction for easy electron-hole separation [605, 997, 1000, 1006]. In a TiO₂-ZnO composite, the electron transfer occurs from the conduction band of ZnO to TiO₂, and the holes move from the valence band of TiO₂ to ZnO [605, 1005] (Fig. 48f). When processing TiO₂-ZnO composite by HPT, the stabilization of high-pressure phases generates additional interphase boundaries between phases which positively contribute to the increment of active sites [605]. Therefore, SPD can contribute to the enhancement of photocatalytic hydrogen production through different routes such as the stabilization of vacancies, high-pressure phase transformations and heterojunction formation.

## 6.6. Electrocatalytic hydrogen production

As mentioned in an earlier chapter, hydrogen is emerging as a new

energy carrier and is gaining global acceptance as a potential fuel pathway. Unlike all carbon-based fuels, hydrogen can be truly carbon-neutral [1007–1009]. Electrocatalytic water splitting to produce renewable hydrogen has been widely regarded as a promising strategy for the coming hydrogen society [1010, 1011]. Similar to photocatalytic water splitting discussed earlier, the electrocatalytic water-splitting process is also composed of two half-reactions [1012]. The oxygen evolution reaction (OER) happens at the anode to produce O₂, while the hydrogen evolution reaction (HER) happens at the cathode to produce hydrogen. Equation 13 is the reaction of water splitting which can be separated into two half-cell reactions (Equations 14 and 15).

$$H_2O \rightarrow H_2 + \frac{1}{2} O_2 \quad E^\circ = 1.23V \text{ vs. SHE} \quad (13)$$

$$2H^+ + 2e^- \rightarrow H_2 \quad E^\circ = 0V \quad (14)$$

$$2H_2O \rightarrow O_2 + 4H^+ + 4e^- \quad E^\circ = 1.23V \quad (15)$$

Developing efficient and stable electrocatalysts for HER and OER of water splitting is very important for electrocatalytic hydrogen production [1013, 1014]. Noble metal catalysts, such as Pt-, Pd- and Ru-based compounds have shown excellent electrocatalytic water splitting catalytic performance [1015, 1016], however, their application is limited due to high price, few resources, etc. Amorphous alloys, which can be prepared by rapid solidification [1017, 1018], have shown promising catalytic and electrocatalytic properties because of the long-range disordered atomic structure [1019–1022]. Moreover, it was reported that amorphous alloys can show significantly enhanced HER and OER performances of water splitting after being processed by SPD [845]. Fig. 49a shows the schematic illustration for the fabrication of Fe₇₈Si₉B₁₃ nanoglass alloy by melt spinning followed by HPT [1023]. Fig. 49b shows that the nanoglass requires an overpotential of only 112 mV to achieve the current density of 10 mA cm⁻² in 1.0 M KOH, which is much smaller than those of the corresponding amorphous alloy (385 mV) and crystalline alloy (465 mV). Similar improvements in HER properties by HPT can also be seen in the Fe₇₃.₅Si₁₃.₅B₉Cu₁Nb₃ [1024] and Fe₇₅B₂₅ [1025] amorphous alloy.

Since the Fe-based amorphous alloys are usually quite hard, the HPT treatment would induce nanocrystallization upon the amorphous matrix, and the effect of nanocrystals on the catalytic performance is difficult to elucidate. A Pd₄₀Cu₃₀Ni₁₀P₂₀ amorphous alloy, which is much softer than Fe-based alloys, is subjected to HPT treatment to study the HER catalytic performances. The HPT treatment did not cause crystallization. The overpotentials at 10 mA cm⁻² of the HPT-treated Pd₄₀Cu₃₀Ni₁₀P₂₀ amorphous alloys were 76 mV and 209 mV in 0.5 M H₂SO₄ and 1.0 M KOH, respectively, which were much smaller than those of 179 mV and 379 mV for the melt-spun Pd₄₀Cu₃₀Ni₁₀P₂₀ amorphous alloy at the same conditions [1015]. The improved HER performances should be mainly attributed to the significantly increased density of flow units in the amorphous matrix by SPD [1027].

FeOOH nanocrystals have been explored as remarkable OER catalysts in past years [1028, 1029]. Based on a previous study, the HPT-treated Fe₇₈Si₉B₁₃ nanoglass with an *in situ* formed FeOOH layer on the surface was prepared by a cathodic corrosion activation method. The FeOOH-covered nanoglass showed excellent OER performances because the FeOOH layer not only exposes more active sites but also generates hydroxyl radicals (OH⁻) in oxygen vacancies to activate the Fe-based nanoglass matrix. Fig. 49c shows that the FeOOH-covered nanoglass generates an overpotential of only 240 mV at 10 mA cm⁻¹ in 1 M KOH, and the Tafel slope is as low as 42 mV dec⁻¹ [1026]. The OER catalytic mechanism has been revealed as shown in Fig. 49d. Firstly, under the action of an electric field, electrons are separated from oxygen vacancies to form positively charged oxygen vacancies (Vo⁺) [1030].

$$Vo \rightarrow e^{-} + Vo^{+} \quad (16)$$

Secondly, Vo⁺ adsorbs the hydroxide ions and water on the catalyst
<figure><img src="image_122.png" /><figcaption>Fig. 49. (a) Schematic illustration for the fabrication of Fe<sub>78</sub>Si<sub>9</sub>B<sub>13</sub> nanoglass alloy [1023]. (b) Linear sweep voltammetry (LSV) curves of the Fe<sub>78</sub>Si<sub>9</sub>B<sub>13</sub> amorphous alloy upon HPT treatment [1023]. (c) LSV curves (iR compensation 90%) of melt-spinning glass (MSG), HPT-processed nanoglass (NG), glass covered by FeOOH (FeOOH@MSG) and nanoglass covered by FeOOH (FeOOH@NG) [1026]. (d) OER catalytic mechanism of the FeOOH@NG [1026]. (e) stability performance obtained from chronoamperometric curves of all samples at the working voltage at 10 mA cm<sup>-2</sup> [1026].</figcaption></figure>

surface to produce hydroxyl radicals [1031, 1032].

$$ \mathrm{OH}^{-} + \mathrm{Vo}^{+} \rightarrow \mathrm{OH}^{+} + \mathrm{Vo} \qquad (17) $$

$$ \mathrm{H}_{2}\mathrm{O} + \mathrm{Vo}^{+} \rightarrow \mathrm{H}^{+} + \mathrm{OH}^{+} + \mathrm{Vo} \qquad (18) $$

Thirdly, hydroxyl radicals attack Fe (0) to produce Fe (II).

$$ \mathrm{Fe}(0) + 2\mathrm{OH}^{-} \rightarrow \mathrm{Fe}(\mathrm{II})(\mathrm{OH})_{2} \qquad (19) $$

Afterward, Fe (II) continues to be oxidized to produce oxygen molecules [1031, 1033, 1034]:

$$ \mathrm{Fe}(\mathrm{II})(\mathrm{OH})_{2} + \mathrm{OH}^{-} \rightarrow \mathrm{Fe}(\mathrm{II}\&\mathrm{III})(\mathrm{OOH}) + \mathrm{H}_{2}\mathrm{O} \qquad (20) $$

$$ \mathrm{Fe}(\mathrm{II}\&\mathrm{III})(\mathrm{OOH}) + \mathrm{OH}^{-} \rightarrow \mathrm{Fe}(\mathrm{III})(\mathrm{ads}) + \mathrm{O}_{2} + \mathrm{H}_{2}\mathrm{O}(21) $$

Finally, the reaction products Fe (III) near the amorphous matrix continuously react with Fe (0) to produce intermediate product Fe(II) (OH)<sub>2</sub>, which participates in the catalytic cycle process [1035, 1036].

$$ 2\mathrm{Fe}(\mathrm{III})(\mathrm{ads}) + \mathrm{Fe}(0) \rightarrow 3\mathrm{Fe}(\mathrm{II})(\mathrm{ads}) \qquad (22) $$

$$ \mathrm{Fe}(\mathrm{II})(\mathrm{ads}) + 2\mathrm{OH}^{-} \rightarrow \mathrm{Fe}(\mathrm{II})(\mathrm{OH})_{2} \qquad (23) $$

In addition, Fe (III) eventually exists in the alkaline medium as Fe (III) (OH)<sub>3</sub> colloids in the absence of Fe (0) by the following reaction.

$$ \mathrm{Fe}(\mathrm{III})(\mathrm{ads}) + 3\mathrm{OH}^{-} \rightarrow \mathrm{Fe}(\mathrm{III})(\mathrm{OH})_{3} \qquad (24) $$

Furthermore, the catalytic performance of the FeOOH-covered nanoglass remains almost unchanged after 50 h continuous reaction (Fig. 49e). In short, the FeOOH-covered nanoglass can perfectly combine self-activation and synergistic catalysis between volume and surface, leading to high activity, fast reaction kinetics and stability of activity for OER. Besides good electrocatalytic activity for hydrogen production, SPD-processed nanoglasses are expected to show good activity for hydrogen storage [1037].

In summary, SPD can greatly enhance electrochemical hydrogen production through the formation of new nanoglass structures. After SPD, the overpotentials at the current density of the amorphous alloys can be greatly reduced, and the stability of HER and OER can be greatly enhanced. It can be predicted that SPD-processed materials will make outstanding contributions to electrochemical catalytic applications.

## 6.7. Hydrolytic hydrogen production

In recent decades, hydrogen generation from hydrolysis of some special materials has been an important research topic [1038]. Fig. 50a shows the commonly used materials for hydrolytic hydrogen production. Among them, active metals like magnesium and aluminum can react chemically with water with the following fundamental reaction equation:
<figure><img src="image_123.png" /><figcaption>Fig. 50. (a) Materials used for hydrolytic hydrogen production [1038]. (b) Hydrolytic hydrogen flow rate curves of Al - 5 wt% LiH - x wt% Ga₂O₃ (x = 1, 3, 5, 7 and 10) at 298 K [1039]. (c) Hydrolytic hydrogen production curves of HPT-processed Al-30Bi-10 C (wt%) composite with different HPT turns [1040]. (d) Combination of three strategies to promote aluminum hydrolysis by high-pressure torsion [1040].</figcaption></figure>

$$M + nH_2O \rightarrow M(OH)_n + n/2 H_2 \quad (25)$$

where M represents the metal and n stands for the number of electrons transferred in the reaction [1038, 1041]. This hydrolysis reaction, however, is often inhibited by the oxide film on the metal surface. Adding some catalysts or destroying the oxide film via mechanical activation are basic strategies to enhance the hydrogen generation behavior (Fig. 50b) [1038, 1041–1043]. SPD is one of the effective methods that can combine the advantages of the above strategies [846]. Successful attempts have been made in the Al-Sn and Al-Bi systems via HPT [1040, 1044, 1045]. In the Al-Sn binary system [1044], the improved hydrogen generation behavior was mainly attributed to the increased amount of micro-galvanic cells between aluminum and tin after HPT processing. The further addition of zinc to the Al-Sn system can promote pitting corrosion and therefore enhance hydrogen production [1045]. As bismuth can play both the role of tin and zinc, binary Al-Bi alloys show much higher hydrogen yield and hydrogen generation rate than Al-Sn alloys [1040]. The addition of graphite to the Al-Bi alloys can further enhance the hydrolysis speed due to the easy fragmentation of the reaction products (Figs. 50b and 50c).

Currently, most research on hydrogen production from the hydrolysis of Al-based materials is focused on finding materials with high reaction rates. For example, the instantaneous hydrogen flow rate of Al-5LiH-5Ga₂O₃ (wt%) is as high as 3837 mL H₂ min⁻¹ g⁻¹ (Fig. 50b) [1039]. Although the high rate is important for emergency cases, a stable hydrogen flow rate is more favorable for fuel cells [1046]. An ideal material may be one that can meet these two directions of application with proper treatments. As shown in Fig. 50c, HPT-processed Al-30Bi-10 C (wt%) composite can exhibit high instantaneous (10 HPT turns with a maximum hydrogen flow rate of ~1600 mL min⁻¹ g⁻¹) or stable hydrogen flow rate (5 HPT turns with a stable hydrogen flow rate of ~600 mL min⁻¹ g⁻¹) depending on HPT turns [1040]. However, the

underlying kinetic as well as the catalytic mechanism of such materials remain complex and need more in-depth studies. For kinetic studies, the shrinking core model is currently a common method to predict and simulate the practical experimental data of the aluminum hydrolysis reaction [1047–1050]. With proper modifications (e.g. considering the impact of single product ALOOH densification with time), this simulating model can correlate well with the actual data [1051–1053]. However, the aqueous solutions of such studies are primarily using acidic or alkaline solutions which is not eco-friendly and few studies have been done on the hydrolysis of bulk materials in neutral solution. Besides the common catalytic effect of catalysts, the catalytic mechanism for the hydrolysis of severely plastic-deformed Al-based materials can be influenced by deformation-induced microstructures. The accelerating effect of micro/nano galvanic cells on aluminum hydrolysis can be mainly attributed to the reduced grain size [1054, 1055]. In addition to the grain size effect, high-angle grain boundaries as well as the deformation texture may also have positive effects on the hydrolysis reaction [1056, 1057]. Several theoretical calculations also pointed out that aluminum hydrolysis can also be improved by some special surface orientations [1058–1060].

As for Mg-based materials, the hydrolysis reaction is similar to the Al-based materials but often occurs at higher temperatures and the passive oxide film is also the key problem. The methods for eliminating the passivation layer include adding acid [1061], reducing particle size [1062], alloying [1063, 1064] and adding strong acids and weak alkaline salts [1065, 1066]. Adding acid can effectively remove the passivation layer, but acid may change the hydrolysis reaction path, which will greatly reduce the theoretical hydrogen yield. Reducing the particle size of magnesium can improve the reaction conversion rate, but too small a particle size increases the difficulty of storage. Adding strong acids and
weak alkaline salts to promote magnesium hydrolysis is a practical and
very effective way [1067, 1068]. With the combination of the above
strategies, MgLi-10EG (wt%) alloy exhibited a hydrogen yield of ~100%
and a maximum hydrogen generation rate of 1147 mL min⁻¹ g⁻¹
[1069].

In summary, SPD can combine the advantages of several effective
strategies for hydrolytic hydrogen production such as the nano-galvanic
cell effect, pitting corrosion effect and fragmentation effect, as shown in
Fig. 50d. Future work should pay more attention to the kinetic studies
and catalytic mechanism of materials with deformation-induced

<figure><img src="image_124.png" /><figcaption>**Fig. 51.** (a) Structural and geometrical features relevant to hydrogen sorption in nanomaterials [1081]. The growth of the hydride phase, depicted in green, requires H₂ dissociation possibly favored by a catalyst, subsurface penetration and H diffusion, which is accelerated along internal interfaces or line defects (dislocations). (b) Dehydrogenation kinetic measurements obtained at 573 K and 1 kPa for nanocrystalline magnesium powders catalyzed by Nb₂O₅ and/or carbon nanotube (CNT) and for the corresponding HPT-processed discs [1086]. (c) High-resolution lattice image of cycled HPT-processed Mg + Nb₂O₅ + CNT composite (inset: selected area electron diffraction pattern) [1086]. (d) Hydrogen desorption curves of cold rolled AZ91 alloy at room temperature (CR) and at low temperature (LTR) [1098]. (e) XRD pattern of a Mg-22Ni (wt%) alloy after fast forging at 480°C with the significant formation of Mg₂Ni [1102]. (f) Hydrogen absorption kinetic curves of an ECAP and ECAP + ARB processed ZK60 alloy measured at 623 K, under 2 MPa pressure [1103].</figcaption></figure>
microstructures.

## 6.8. Solid-state hydrogen storage

Hydrogen-based energy storage can be a promising technology shortly [1070, 1071]; however, there are still technological challenges in the way of realizing it in practice. One obstacle is the difficulty of storing hydrogen. A hydrogen storage system has to possess high storage capacity, appropriate charge/discharge rate and long operational cycle life [1072]. In the past decades, a lot of effort was put into the development of material-based hydrogen storage systems, including metal hydrides [1073, 1074], as they generally offer large storage capacity by volume and reversible operation [1075, 1076]. Metal hydrides store hydrogen not only on the surface but in the whole volume of the material as well through a chemisorption reaction [1077–1079]. Thus, hydrogenation/dehydrogenation of these materials involves both long-range diffusion of hydrogen atoms and phase transformations. Structural inhomogeneities, nanostructuring and lattice defects can promote these processes [1080]. It was reported that altering the microstructure by SPD methods is an effective way to enhance the hydrogen absorption and desorption kinetics of these materials, see Fig. 51a [847, 1081, 1082]. Moreover, mechanical treatment by SPD can induce the formation of dislocations and grain boundaries in the oxide layer that is usually present on the surface of metal hydrides, thus making it permeable to hydrogen for easy activation. In this section, various SPD methods applied to hydrogen storage materials are reviewed and the impact of such process on the thermodynamics, kinetics and activation of these materials will be discussed in coming sections.

* HPT can be applied to various hydrogen storage systems, as demonstrated in a recent overview [358]. When MgH₂ powder is deformed by HPT, an intensive nanostructure formation takes place, which positively influences the hydrogen storage performance [1083]. The application of high-energy ball milling before HPT can promote mechanochemical routes, resulting in the formation of intermetallic compounds with advantageous hydrogen storage properties. Plastic strain generated by HPT can increase the maximum hydrogen capacity of ball-milled Mg₇₀Ni₃₀ nanopowders by 50 % due to an increased number of defects [1084]. It was established recently that carbon nanotubes (CNTs) can provide fast diffusion channels during hydrogenation and dehydrogenation for the hydrogen atoms through the passivated surface layer into the bulk material [1085]. The addition of Nb₂O₅ to magnesium catalyzed by CNTs plays a crucial role in attaining suitable hydrogen capacity (see Fig. 51b) [1086]. The CNT sections presented in Fig. 51c are preserved during the plastic deformation of ball milling + HPT and subsequent sorption cycling. The hydrogenation behavior of fully disordered systems, like metallic glasses can significantly be improved when the material is subjected to severe shear deformation by HPT [1087, 1088]. The enhanced hydrogen sorption kinetics of a melt-spun Mg₆₅Ni₂₀Cu₅Ce₁₀ nanoglass subjected to torsional straining are attributed to the abundant pathways for hydrogen diffusion at the interfaces between the nanoglass regions that are developed during the HPT process [1037].

* ECAP is often used to enhance the hydrogen storage properties of different bulk materials. It was shown that the number of ECAP passes [1089] and the deformation route [1090] have a significant effect on the absorption/desorption time and the maximum reversible capacity of different Mg-based alloys. The improvements were attributed to the increased grain boundary fraction that presents nucleation sites for the hydride and provides fast diffusion channels for hydrogen atoms. The use of ECAP can also induce the crystallization of different second-phase particles (for example, Mg₁₇Al₁₂ in AZ61 and AZ31 alloys) that can enhance the absorption/desorption performance of the material as catalysts [1091, 1092]. Improvement

of the long-term cycling stability up to 1000 cycles was demonstrated for ZK60 magnesium alloy after ECAP processing [1093].

* Cold rolling was shown to reduce the necessary activation time for different Mg-based materials [1094, 1095] and the TiFe alloy [1096, 1097]. Rolling can also be performed at cryogenic temperatures, which results in a more refined microstructure and better absorption/desorption kinetics than the sample rolled at room temperature (see Fig. 51d) [1098]. ARB offers the possibility to synthesize composites in which the secondary phase particles can act as nucleation sites as was demonstrated for the Mg-Ti system [1099]. It was shown that cold rolling induces a strong (002) texture in magnesium and this texture promotes the hydrogen absorption capability of the specimen.

* Fast forging, a less conventional SPD technique, can operate at larger scales and shorter times [1100]. Earlier research on the mechanical processing of MgH₂ + 5 wt% Fe by fast forging resulted in a good dispersion of the catalyst particles and significant nanostructuring [1101]. A Mg+Ni powder mixture processed by fast forging exhibits a threshold processing temperature, below which hydrogen absorption is mainly governed by the formation of defects and cracks, enhancing hydrogen diffusion into the bulk, however, forging at elevated temperature yields the solid-state formation of Mg₂Ni nanocrystals that can act as a catalyst, see Fig. 51e [1102].

* Combination of different SPD techniques, for example, ARB + ECAP [1103] and ECAP + cold rolling [1104], can further improve hydrogen storage performance compared to the use of only a single method (see Fig. 51f).

In conclusion, several SPD methods can significantly improve the hydrogen storage capabilities of different classes of materials via refining their microstructure, generating grain boundaries, and inducing lattice defects, hence improving the diffusion of hydrogen and the nucleation of the hydride phase. Nevertheless, the different SPD techniques promote different microstructures (different types and densities of defects, texture, etc.) which can result in different hydrogen sorption properties.

### 6.8.1. Thermodynamic tailoring of hydrogen storage materials

Solid hydrogen storage materials such as solid-state hydrides have been studied for over 50 years [1105], for their uses as energy storage materials [1106–1108]. As discussed in an earlier chapter, the merits of these materials include lower operation pressure and energy consumption, higher volumetric hydrogen density and safety compared with compressed high-pressure or cryogenic liquid hydrogen storage [1109–1116]. A main concern regarding these materials is tailoring their thermodynamics. Hydrogen is reversibly absorbed and desorbed via a chemical reaction accompanied by energy change [1117].

$$M + x/2H_2 \rightleftharpoons MH_x + \Delta H \quad (26)$$

where M is the metal or alloy, x is the amount of reacted hydrogen and ΔH is the enthalpy change. Fig. 52a illustrates the reversible phase transition process between metal (in the α phase) and metal hydrides (in the β phase) [1118]. Hydrides can be divided into two categories of interstitial and non-interstitial hydrides, which are comprised of metallic and covalent or ionic bonds, respectively [1119]. The dehydrogenation temperatures of solid-state hydrides are positively correlated with bond strength [1120]. This means the dehydrogenation temperature of the hydrides is determined by thermodynamics, which can be obtained from the van't Hoff equation [1121, 1122]:

$$\ln PH_2 = -\frac{\Delta H}{RT} + \frac{\Delta S}{R} \quad (27)$$

In Equation 27, P is hydrogen pressure, T is temperature, R is gas constant and ΔS is the entropy change which is approximately equal to 130 J/(K mol H₂) [1123, 1124]. Fig. 52b indicates the corresponding
<figure><img src="image_125.png" /><figcaption>Fig. 52. (a) Diagrammatic representation of phase transition during reversible hydrogen storage process [1118]. (b) Typical pressure-composition isotherm, phase diagram, and van't Hoff plot of a metal hydride [1122, 1123].</figcaption></figure>

<figure><img src="image_126.png" /><figcaption>Fig. 53. (a) XRD pattern of Mg₄NiPd after 1500 HPT turns [359]. (b) Pressure-composition isotherm of Mg₄NiPd at 305 K [359]. (c) Pressure-temperature phase diagram of MgH₂ [515]. (d) Heat flow in differential scanning calorimetry for MgH₂ processed by HPT for various turns [515].</figcaption></figure>

relationship between pressure-composition-temperature (PCT) iso-
therms and the van't Hoff equation [1122, 1123]. For the reversible
desorption/absorption of hydrogen at a temperature close to room

temperature, the ideal dehydrogenation enthalpies of hydrides are
25–35 kJ/mol H₂ [1124, 1125]. Therefore, thermodynamic tailoring is
in great need to design materials for practical applications [1126].
Thermodynamic tailoring can be realized by the following approaches: (i) nanoengineering [1127–1129]; (ii) alloying [1118, 1129–1132]; (iii) phase transition control [1101, 1133]; and (iv) reaction pathway changes [1129, 1134, 1135].

SPD has been proven to be a powerful tool for thermodynamics tailoring, by taking MgH₂ as an example. MgH₂ is widely regarded as a promising hydrogen storage material because (i) its high hydrogen capacity is 7.6 wt%, and (ii) magnesium is the ninth most abundant element in the universe [1129]. High dehydrogenation thermodynamics of 75 kJ/mol H₂ for MgH₂, however, hinders its application at room temperature [1129]. Thermodynamic destabilization of magnesium is a main topic for hydrogen storage at room temperature. Recently, a Mg-based ternary alloy Mg₄NiPd was designed based on the prediction of first-principles calculations. The Mg₄NiPd alloy was expected to have moderate hydrogen binding energy between the Mg₂Ni and Mg₂Pd ones [359, 1077]. It was successfully prepared with an ordered BCC-based structure (Fig. 52a) through 1500 turns of HPT processing and was found to exhibit a reversible 0.7 wt% hydrogen storage capacity at room temperature (Fig. 53a). Among the possible octahedral sites of 4Mg-2Ni, 4Mg-1Ni-1 Pd and 4Mg-2 Pd in Mg₄NiPd, the 4Mg-1Ni-1 Pd site with the hydrogen binding energy of -0.12 eV was found to be responsible for the room-temperature hydrogen storage. Furthermore, a series of alloys with elevated hydrogen storage properties have been fabricated by the HPT process, such as Mg₆₅Ni₂₀Cu₅Y₁₀ alloy [1088], Mg₆₅Ce₁₀. Ni₂₀Cu₅ nanoglass [1037], MgTiVCrFe HEA [314], etc. [1082]. HPT has become a popular technology to improve the performances of many hydrogen storage materials including, but not limited to TiFe alloys [664, 1136, 1137], Mg-based alloys [313, 515, 1083, 1138, 1139], HEAs [1082] and others [192, 309, 322].

Another strategy to tailor the thermodynamics of hydrogen storage materials is through phase transformations. For example, MgH₂ has an α phase with a stable tetragonal structure at ambient pressure and transfers into the γ phase or the β phase at higher pressures (Fig. 53c) [1077, 1133]. A transition from the α phase to the γ phase occurs by HPT processing under 5 GPa, and the proportion of the γ phase increases with the increase of HPT turns. The formation of almost 100% γ phase was obtained after HPT processing for 15 turns. With the increase of the γ phase fraction, the dehydrogenation temperature reduced gradually. The formation of the nanocrystalline γ phase by 15 turns of HPT processing resulted in a decrease of the dehydrogenation temperature by 80 K (Fig. 53d) [515]. The results indicate that HPT processing is also effective in reducing the thermodynamic stability of MgH₂.

In conclusion, SPD is a powerful tool that helps to tailor the thermodynamics of hydrogen storage materials by fabricating metastable alloys or preparing unstable phases. Among the different SPD methods, HPT is currently the most popular method for tailoring the thermodynamics of such materials. However, HPT has main drawbacks to be overcome [1077, 1082]: it is hard to prepare samples on a large scale. In the future, HPT technology and other SPD methods are expected to be further improved for the preparation of samples with not only larger sizes but also improved properties.

### 6.8.2. Kinetic tailoring of hydrogen storage materials

Kinetics of hydrogen storage in metal hydrides, particularly in magnesium and magnesium-based hydrides, are actively investigated because of their potential for commercial applications [1140]. One way to increase the hydrogenation kinetics in metal hydrides is by mechanical deformation. Many techniques have been used, such as ECAP [1093, 1141–1144], HPT [181, 192, 309, 334, 359, 1139, 1145, 1146] and cold rolling (also called ARB) [5, 1147–1150]. This last technique is arguably the easiest to scale up to the industrial level. However, cold rolling is usually performed in a horizontal fashion which may be problematic for metal hydride materials because they are often in a powder form. The solution to this problem is to perform the rolling vertically. In the cold rolling process, the deformation is mostly plane strain compression. The strain tensor is expressed as $ε_{ij}$ (i, j = 1, 2, 3) where 1 is the rolling

direction (RD), 2 is the transverse direction (TD) and 3 is the normal direction (ND). In the ideal case, the strain in cold rolling is $ε_{11} = -ε_{33}$ and all other strains are zeros ($ε_{22} = ε_{12} = ε_{13} = ε_{23} = 0$). The effective strain is then equal to the von Mises equivalent strain [1151].

$$ε = \sqrt{\frac{2}{3}ε_{11}} \qquad (28)$$

Magnesium is ductile but has only four slip planes [1152]. This means that magnesium will experience important work hardening after only a few rolling passes. For hydrogen storage applications, this is not a problem because the metal hydrides are usually used in the form of powders. Also, upon hydrogenation/dehydrogenation a metal hydride will deprecitate. Therefore, work hardening is acceptable for metal hydrides. Many groups have investigated the effect of cold rolling on pure magnesium and they found that cold rolling improved the hydrogenation kinetics [1098, 1141, 1153–1156]. The same conclusion was reached from investigations on magnesium alloys such as AZ91 [1098], MRI153 [1157], ZK60 [1158] and Elektron 21 [1156]. Other systems such as Mg-Al [1159], Mg-Cu [1160], Mg-Fe [1161], Mg-Ni [1162] and Mg-Pd [1163] have been also investigated. Fig. 54a shows a comparison of the first hydrogenation of magnesium and two magnesium alloys (MRI153 and AZ91D) after cold rolling while Fig. 54b shows the second hydrogenation [1164]. It is clear that cold rolling is beneficial for the hydrogenation kinetics. A more in-depth discussion of all these systems could be found in a recent review paper [1165]. Cold rolling can be also used on hydrogenated powder. Fig. 54c shows the effect of cold rolling and ball milling on MgH₂ powder [1166]. We see that the hydrogenation kinetics is much faster for the mechanically processed samples. Ball milling produced a sample with a higher hydrogen capacity and faster kinetics. However, the difference is that ball-milled powder should always be under an argon atmosphere while cold-rolled samples are processed in the air.

In conclusion, SPD methods are efficient in improving the hydrogenation kinetics of hydrogen storage materials. In this regard, cold rolling of magnesium alloys is promising because of its potential for scaling up and industrial use, although there are still many technical investigations that need to be done. For example, appropriate magnesium compounds should be searched, and the effect of cold rolling parameters such as rolling speed, number of rolls, temperature of rolling, number of rolling passes, thickness reduction at each pass, etc. should be studied. But arguably the most important investigation is to get a more in-depth understanding of the exact mechanism that makes the kinetics of hydrogenation better after SPD processing.

### 6.8.3. Activation and air resistance of hydrogen storage materials

In addition to thermodynamics and kinetics, activation is another important issue that should be considered for the industrial application of hydrogen storage materials. Difficult activation results in no hydrogenation, even when the material is thermodynamically favorable for hydrogen storage. Difficult activation is usually due to the formation of an oxide layer in air on the surface which hinders hydrogen dissociation and diffusion. The formation of oxide reduces the air resistance of these materials, and thus, they should be usually handled under an inert gas atmosphere. Hydrogen storage materials are activated either by heating to high temperatures under vacuum, chemical or mechanical modification. Perhaps, TiFe is the most famous hydrogen storage material that suffers from difficult activation. TiFe has a B2 crystal structure and is known as a good candidate for stationary (housing) and semi-stationary (submarines) hydrogen storage applications because its hydrogenation is thermodynamically possible at room temperature, its volumetric storage capacity is high and it is reasonably cheap [1167–1172]. Nevertheless, the practical use of TiFe is hindered due to the difficulty of initial activation for the hydrogenation which requires exposure to a vacuum or a hydrogen atmosphere under a pressure of ~3 MPa at temperatures higher than ~673 K [1167–1172]. SPD has introduced a
<figure><img src="image_127.png" /></figure>

<figure><img src="image_128.png" /></figure>

<figure><img src="image_129.png" /><figcaption>**Fig. 54.** (a) First and (b) second hydrogenation at 623 K under 2 MPa of hydrogen for cold-rolled pure magnesium, MRI153 and Az91D alloys [1157, 1164]. (c) Normalized kinetic absorption at 623 K and under 2.0 MPa hydrogen pressure for MgH₂ in as-received, cold rolled (five passes, CR5X) and mechanically milled for 30 min (BM 30 m) [1166].</figcaption></figure>

new pathway for easy activation of TiFe and many other materials as will be discussed in this section [664, 1136, 1137, 1173].

An example of the effect of SPD on activation is shown in Fig. 5a where TiFe was processed by HPT for different numbers of turns including the samples after annealing and after compression but without rotation [1136]. While hydrogen absorption is essentially negligible in

<figure><img src="image_130.png" /></figure>

<figure><img src="image_131.png" /><figcaption>**Fig. 55.** (a) pressure-composition isotherms of TiFe at 303 K for annealed sample, sample processed by compression under 6 GPa (N = 0) and samples processed by HPT for N = 1/4–10 turns [1136]. (b) Variation of hydrogen content against hydrogenation time at room temperature after the first and second exposure to hydrogen under a pressure of 2 MPa for Ti-Fe-Mn intermetallic processed by HPS including as-received ingot and HPT-processed disc [1186].</figcaption></figure>

the samples after annealing and after compression, the absorption occurs by 1.6–2.1 wt% after HPT processing. It is important to note that this hydrogen absorption was realized after exposure to air for several hundred days following HPT processing. Moreover, when an HPT-processed sample was left in the air for 400 days, subjected to hydrogenation for five cycles, and left in the air for an extra 30 days, hydrogen absorption still occurred without activation. Such experiments thus demonstrate that HPT-processed TiFe is quite air resistant and hardly deactivated in the air once it is processed by HPT [1136]. Here it should be noted that TiFe is an extremely hard and brittle material, but its SPD processing could be realized by HPT [1, 2, 23, 177, 179]. As discussed earlier, the high processing pressure in HPT makes it applicable not only to metallic materials [177, 674, 1174] but also to hard-to-deform materials such as intermetallics [1175–1178], ceramics [584, 844, 1179] and semiconductors [512, 1180–1185].

Activation of hydrogen storage materials by SPD is not limited to the HPT method. As an alternative to the HPT process, which uses small disc- or ring-shapes, the HPS process has been developed [57]. The HPS utilizes sheets, and thus, it has the potential to upscale the sample size [179, 240, 1187]. As mentioned earlier, the strain in HPS is accumulated by the reciprocation of anvils with respect to the sample [1188], while the rotation of the anvils is adopted in HPT [20]. Because HPS is operated under high hydrostatic pressure as in HPT, its application is also
feasible to hard-to-deform materials [240, 244, 1187]. In a recent study, a TiFe intermetallic containing a minor addition of Mn (TiFe<sub>1-x</sub>Mn<sub>x</sub> with x = 0.3) was processed by HPS to enhance its hydrogen storage activation [1186]. The addition of Mn to TiFe reduces the hydrogenation pressure and activation temperature [1189, 1190], but the material still needs activation [1173]. Fig. 55b shows the hydrogenation kinetics curves of the HPS-processed strip and an as-received ingot [1186]. For comparison, a hydrogenation curve is also included from the HPT-processed sample [1173]. The hydrogenation occurs quickly in the sample processed by HPS, while no appreciable hydrogenation proceeds in the as-received ingot without the HPS process, at least for the period exposed to the hydrogen atmosphere. Incubation of ~60 min is observed after the first exposure of the HPS-processed sample to hydrogen, but the hydrogenation starts immediately after the second exposure in the HPS-processed sample. The fast hydrogenation for the second exposure without incubation is similar to the one observed in the HPT-processed sample (measured after the fourth hydrogenation cycle) [1173].

The activation of hydrogen storage materials by SPD processing was also reported by cold rolling [1137], rolling and HPT [1191] and ball milling [1192, 1193]. Here, it should be noted that the materials processed by ball milling as a powder-SPD method are not air resistant, while materials processed by cold rolling, HPS and HPT as bulk-SPD methods are quite air resistant. As discussed earlier [1136], the enhanced activation and air resistance by bulk-SPD methods should be due to the formation of a high density of lattice defects such as dislocations and grain boundaries. Particularly, grain boundaries are expected to provide pathways for fast hydrogen diffusion and improve both hydrogenation activation and kinetics [1136]. Future studies are needed to clarify the exact mechanism underlying the positive effect of SPD on hydrogen storage activation.

### 6.8.4. Mass sample production for hydrogen storage

Magnesium is one of the most appropriate hydrogen storage materials [1194, 1195], while SPD-processed magnesium has both fast kinetics and large hydrogen uptake, as discussed in recent review papers [5, 1077, 1196]. Hydrogen diffusion, restricted solid solution of hydrogen in magnesium and nucleation of hydrides are some issues that are positively affected by SPD processing [1197–1200]. However, mass production of SPD-processed materials is a main challenge for their commercialization in hydrogen storage applications. Besides cold rolling, as reviewed in an earlier chapter, fast forging (high-speed drop forging) and ECAP are two methods with high potential for upscaling, as will be discussed in this section.

The so-called fast forging process (Fig. 56a) was applied to Mg-based materials. The flow of material takes a rather perfect two-dimensional isotropic character in the open die fast forging procedure contrarily to the one-dimensional transformation of the microstructure in cold rolling. Moreover, compared to cold rolling with a 50% thickness reduction, a much higher thickness reduction (90%) can be achieved by fast forging in only one pass [1201]. For example, a cylinder billet of a Mg-rich alloy with 12 mm in diameter and 20 mm in height, transforms into a disc of ~38 mm in diameter and ~2 mm in height (thickness) in one pass of fast forging [1202]. Another benefit of fast forging is that the process is operated fast within $5 \times 10^{-3}$ s with a flow rate of at least 25 m/s [1100, 1203]. In fast forging, the kinetic energy E of the free-falling load is assumed to expand mostly on the impact energy of plastic deformation $\Delta W$ [1204]. The heat flux of the adiabatic process on impact takes the form $\rho c \Delta T \sim k_p \Delta W$ ($\rho$: density, c: heat capacity, T: Temperature $k_p$: the coefficient determining the fraction of sample converted by plastic deformation at a given internal heat level). Other specific parameters like energy loss due to external friction with both the hammer and the die, intense friction between grains, formation of cracks and elastic vibrations can be defined by an additional k coefficient, leading to $\Delta T \sim k_k \Delta W/\rho c \sim 200$ K. This thermal effect in fast forging enables parts of the sample to reach an intermediate situation between fragile to ductile states [1205, 1206]. It can be shown using the Crockroft and Latham

<figure><img src="image_132.png" /></figure>

<figure><img src="image_133.png" /><figcaption>**Fig. 56.** (a) Fast Forging instrument of the drop type (sample is placed in an argon-filled chamber and can be also heated using an induction coil). (b) ECAP instrument (die can be heated on demand by resistive inserts placed in the die. The deformation rate can be controlled in both instruments by computer-connected sensors [1100].</figcaption></figure>

criterion [1207, 1208] that the thermal effect and non-uniform metal flow may generate porosities, de-cohesion then create extended fractures which are positive features for hydrogen storage materials.

Since the early paper published by Skripnyuk *et al.* [1209], ECAP has been demonstrated as an efficient SPD method for the mass production of Mg-based materials with fast hydrogen sorption kinetic [1045, 1143, 1210–1213]. After most of ECAP procedures, high-energy ball milling is applied to hydrogen storage materials to obtain a higher specific surface, but the time for milling is quite short [1214]. Nevertheless, the ECAP route showed high potential for the delivery of various hydrogen-reactive Mg-based materials such as ZK60 [1103], AZ31 [1215]) or more specific alloys [1216]. However, the sorption kinetics after ECAP are not still as fast as the widely reported data reported in the literature for the ball-milled powders with catalyst addition [1217–1223]. An ECAP instrument used to process hydrogen storage materials at the laboratory scale is shown in Fig. 56.

The performance of fast forging and ECAP for hydrogen storage materials is reasonably comparable notwithstanding the study where powders of pure magnesium (5–20 µm grain size) were compacted and processed by both fast forging and two passes of ECAP [1224]. The mechanically treated materials remained particularly fragile, so no subsequent high-energy ball milling was needed, but only attrition was applied for a few tens of seconds. Table 6 summarizes the results of hydrogen absorption and desorption, indicating the performances of the two methods are reasonably comparable.
<table><caption>**Table 6**
Hydrogen absorption and times and temperature for magnesium powder processed by ECAP and fast forging [1224].</caption><thead><tr><th>SPD Process</th><th>Absorption (~7 w %)</th><th>Time (min.)</th><th>Desorption (100 w %)</th><th>Time (min.)</th></tr></thead><tbody><tr><td>ECAP (2B<sub>c</sub>)</td><td>2 MPa, 380°C (cycle 3)</td><td>45</td><td>15 kPa, 380°C (cycle 3)</td><td>30</td></tr><tr><td>Fast Forging</td><td>2 MPa, 380°C (Cycle 1)</td><td>60</td><td>15 kPa, 320°C (cycle 1)</td><td>60</td></tr></tbody></table>

The reduction of the grain sizes to the UFG level (multiplication of the grains boundaries) and simultaneous creation of pores and cracks appear to be desirable to achieve such enhanced properties. The addition of catalysts during mass production to these UFG materials with pores and cracks can provide additional benefits for hydrogen storage [1225]. It is expected that the knowledge produced in recent years for SPD can facilitate the mass production of metal hydrides, particularly using methods such as cold rolling, fast forging and ECAP.

## 6.9. Hydrogen embrittlement resistance

The susceptibility to hydrogen embrittlement is markedly increased in high-strength steels with a tensile strength higher than 1.2 GPa. The addition of carbide-forming elements such as titanium, vanadium and molybdenum and subsequent high-temperature tempering are designed not only to contribute to secondary hardening but also to provide irreversible trapping into the martensite matrix to prevent local hydrogen concentration at prior austenite grain boundaries in BCC steels [1226-1228]. The roles of the lattice defects in the hydrogen embrittlement mechanisms are widely varied with different crystal structures, i.e. BCC and FCC structures [1229]. For austenitic stainless steel with an FCC structure, titanium-bearing Fe-25Ni-15Cr (wt%) steel is recognized as precipitation-hardened austenitic steel, which is strengthened via the dispersion of fine Ni<sub>3</sub>(Al,Ti) precipitates. However, this type of steel is susceptible to hydrogen embrittlement despite containing large amounts of rare metal elements [1230]. Solid solution strengthening by interstitial atoms such as nitrogen enhances hydrogen-induced strain localization, which reduces the hydrogen embrittlement resistance [1231]. By contrast, Macadre *et al.* reported that grain refinement to approximately 1 µm mitigated the hydrogen-induced ductility loss in

Fe-16Cr-10Ni steel [1232]. Further, Mine *et al.* investigated [1233, 1234] the hydrogen effects on the tensile properties of UFG stainless steels produced by HPT processing and demonstrated that post-HPT annealing reduced the hydrogen-induced ductility loss while maintaining their high strength by reducing the dislocation density and mixing secondary phases in UFG austenite. Recently, it has been reported that high strength and reduced hydrogen embrittlement susceptibility were attained by introducing UFG and nanotwins into HEAs with an FCC structure [1235, 1236].

For metastable austenitic steels, although transformation-induced plasticity (TRIP) and twinning-induced plasticity (TWIP) phenomena improve the balance between tensile strength and ductility, the mechanical properties are considerably degraded in the presence of hydrogen [1237, 1238]. Park *et al.* and Bai *et al.* reported that grain refinement suppressed the hydrogen embrittlement in the TWIP steels [1239, 1240]. Metastable austenitic stainless steels such as types 301 and 304 with ordinary grain sizes undergo severe hydrogen embrittlement, which is attributed to deformation-induced martensitic transformation during tensile loading [1241, 1242]. Fig. 57 shows the effect of hydrogen on the deformation and fracture morphology of ordinary-grained type 304 stainless steel [1237]. The uncharged specimen exhibits a typical cup-and-cone failure with dimples on the fracture surface (Fig. 57a). The cross-section area is reduced at a part remote from the fracture surface along with a significant uniform elongation (Fig. 57b). Fig. 57c shows the EBSD phase map of the microstructure near the fracture surface at a longitudinal cross-section of the fractured specimen. Grains are substantially elongated along the tensile axis and α' martensite is densely formed at the fracture surface (Fig. 57c). By contrast, the specimen containing ~25 mass ppm of hydrogen exhibited plastic inhomogeneity and less α' martensitic transformation (Fig. 57d) when compared to the uncharged specimen (Fig. 57c). The hydrogen-charged specimen was fractured on the plane macroscopically perpendicular to the loading axis with many microcracks on the free surfaces (Fig. 57e). In addition, the fracture surface of the hydrogen-charged specimen was composed of quasi-cleavages and flat facets (Fig. 57f). Microtension testing studies successfully characterized the hydrogen-induced quasi-cleavage and flat facet formations using single-crystalline and twin bi-crystalline specimens, respectively [1243, 1244]. Hydrogen-induced quasi-cleavages in metastable austenitic stainless steel occurred at the interphase boundaries between the

<figure><img src="image_134.png" /><figcaption>**Fig. 57.** SEM images of fracture surface, side view and EBSD phase maps at longitudinal cross-section in (a-c) uncharged and (d-f) hydrogen-charged specimens of ordinary-grained type 304 stainless steel [1237].</figcaption></figure>
<figure><img src="image_135.png" /></figure>

<figure><img src="image_136.png" /><figcaption>**Fig. 58.** (a) Grain size dependence of yield stress and (b) relationship between area reduction and ultimate tensile strength for 304 stainless steels with different grain sizes [848].</figcaption></figure>

austenite and α' martensite formed during the tensile loading, which was triggered by excess hydrogen generated owing to the differences in the hydrogen diffusivity and solubility between the two phases [1245]. The flat facet formation associated with twin boundary separation is the primary mechanism for premature failure concerning hydrogen embrittlement of the metastable austenitic stainless steels [1244, 1246]. It is then hypothesized that the dispersion of dynamic martensitic transformation suppresses the local hydrogen concentration at the interphase and twin boundaries.

Fig. 58 shows the effects of hydrogen charge on the tensile properties of type 304 stainless steels with different grain sizes, which is obtained using microtension testing [848, 1247]. Both uncharged and hydrogen-charged specimens well held the Hall-Petch relationship between yield stress and average grain size (Fig. 58a). The friction stress was increased by 40% by hydrogen charging but the Hall-Petch coefficient was invariable. This indicates that hydrogen strengthens the matrix via the solid solution effect, but does not influence the grain

refinement strengthening. The relationship between the reduction of area and UTS in the uncharged specimens reveals that the area reduction is reduced with increasing UTS in the average grain size range larger than ~0.4 µm, whereas it remains unchanged in the smaller grain sizes (Fig. 58b). The hydrogen-charged specimens have 1–1.5 GPa UTS with 60% area reduction in the UFG microstructures with grain sizes smaller than ~0.4 µm. In other words, ultra-grain refinement not only improves the balance between strength and ductility but also mitigates the hydrogen embrittlement susceptibility in metastable austenitic stainless steel. It has been argued that dynamic α' martensitic transformation is the crucial factor determining hydrogen embrittlement susceptibility [1245, 1248, 1249]. The saturated hydrogen concentration is lower in α' martensite than in austenite by an order of magnitude [1250]. When austenite containing hydrogen is transformed to α' martensite, excess hydrogen can be generated, which corresponds to the difference in the saturated hydrogen concentration between the two phases. Subsequently, excess hydrogen diffuses out of the formed α' martensite to the surrounding austenite and is accumulated at the interface boundaries because of the extremely high diffusivity of hydrogen in martensite compared to that in austenite. A finite element calculation study by Wang *et al.* supported [1251] that cracking occurs in the austenite region neighboring the formed α' martensite. UFG austenite formed a single-variant α' martensite from each grain, whereas ordinary-grained austenite selected several α' martensite variants with a favorable orientation [1252]. For this reason, excess hydrogen generated from UFG austenite is dispersed, which results in reduced local hydrogen concentration and in turn, mitigates the hydrogen-induced ductility loss. As shown in Fig. 58b, the area reduction and UTS relationship reveals that the nanotwinned microstructures as well as the UFG microstructures can contribute to the strengthening while reducing the hydrogen embrittlement susceptibility in metastable austenitic stainless steel [848, 1253]. This supported that the hydrogen embrittlement resistance was enhanced in type 304 and TWIP steels with bimodal microstructures consisting of nanotwin bundles and recrystallized grains [1254–1256].

In summary, introducing UFG and nanotwins into metastable austenitic stainless steels is effective for mitigation of hydrogen-induced ductility loss through controlling dynamic α' martensitic transformation. Plastic deformation processes of metastable austenitic steels at controlled moderate temperatures are a promising strategy for obtaining sustainable materials.

## 6.10. Corrosion resistance

With the advent of SPD, the effect of plastic deformation on corrosion has again come back to being an important topic of materials science [849, 1257, 1258]. In the ultrahigh strain range, microstructural features such as dislocation density, misorientation, and structures of grain boundaries develop progressively into an ultimate form of UFG or nanocrystalline structures which can affect the corrosion resistance [454]. For example, the introduction of a very high strain yields a maximum dislocation density of approximately $10^{15} \text{ m}^{-2}$ in pure copper [454, 1259] and iron [446, 1260]. Moreover, the final grain size is copper is reduced to 0.4 µm by ECAP at room temperature [1261, 1262] and 80 nm by HPT at 100 K [1263]. After reaching the complete UFG structure, dislocations inside grains decrease; however, some residual dislocations remain. Such high-density grain boundaries with residual dislocations impact the corrosion behavior even in pure metals or single-phase materials [849]. A similar situation occurs in HEAs or multi-principal elements alloys (MPEAs), which are a new class of materials with superior mechanical properties, and some of them have very high corrosion resistance [1264]. The effect of SPD on the corrosion resistance of metals and conventional alloys is of great importance in scientific and practical viewpoints and the corrosion resistance of SPD-processed HEAs is waiting to be explored [1265]. Although corrosion resistance has been investigated for a wide range of alloys, Fe-based alloys and steel received the most significant attention due to their
<figure><img src="image_137.png" /><figcaption>**Fig. 59.** (a) Hierarchical steps leading to the protective ability of passive films from the corrosive environment and microstructure of base metals [849]. (b) Relation between grain size of UFG stainless steels obtained via various methods versus the difference in pitting potentials of UFG (*E*<sub>b</sub>') and coarse-grained (*E*<sub>b</sub>') materials ("Cr" and "No Cr" indicate that chromium enrichment in passivation films was observed or not) [849]. (c) Relation between the pitting potential of Fe-12Cr (wt%) alloys and grain size after ECAP (eighth passes) and subsequent annealing (the numbers in the figure indicate the annealing temperature) [1275].</figcaption></figure>
practical applications. The formation of a protective “passive film” on the surface and the effect of SPD on this film are the most critical issues studied in recent years. In this section, the features of the passive layer and corrosion resistance of steels are discussed.

### 6.10.1. Corrosion passive film features

Studies have shown that the corrosion resistance of Fe-based alloys in passive environments is enhanced by grain refinement down into the submicron or nanoscale level [1266]. This beneficial effect in Fe-Cr alloys has been attributed to the enhanced protective nature of the passive film due to a greater chromium enrichment in the film [1266]. Two independent mechanisms for greater chromium enrichment in passive films have been proposed: enhanced selective dissolution of iron and faster chromium diffusion [849, 1257, 1266]. Both mechanisms are supposed to originate from high-density of grain boundaries, as reported using high-resolution observation techniques such as high-angle annular dark-field scanning transmission electron microscopy (STEM) [1267] and in situ atomic force microscopy [1268]. Examination of nanocrystalline 316 L stainless steel processed by a surface mechanical rolling treatment (SMRT) has shown that protective passive films with surface chromium enrichment were localized to a zone at grain boundaries with a width of up to 50 nm [1267]. Maurice et al. showed the presence of an inhomogeneous chromium distribution that stems from the grain boundaries of the base metal and the granular structure of the passive film using in situ atomic force microscopy [1268]. This finding suggests that both these mechanisms, facilitated by grain refinement, might be capable of homogeneous passive film formation over the entire surface if the grain size is extremely small (possibly, <100 nm), which cannot be achieved by most classical SPD methods. Therefore, for the formation of a uniform and homogeneous passive film on all grains, the role of factors other than grain size should be considered. Fig. 59a shows the hierarchical relation between the protective ability of passive films formed on base metals of Fe-Cr alloys and structural parameters [849]. The protective ability is affected by the semi-conductive properties such as ion/point defect transport [1269, 1270] and/or chemical stability against dissolution into the solution [1271]. These properties are further related to crystal structure, compactness, structural or chemical inhomogeneity, and chemical composition. The passivity of Fe-Cr alloys appears in neutral to acidic solutions when the bulk chromium content exceeds approximately 12% [1272]. The passivation involves chromium enrichment in the passive film [1273], which has generally been explained based on the selective dissolution of iron and oxidation of chromium at the surface [1274]. Therefore, the structure of passive films is affected by the environment as well as the structure of the base metal.

Fig. 59b depicts the very low correlation between the difference in the pitting potentials of nanocrystals (E<sub>b</sub>’) and their coarse-grained counterpart (E<sub>b</sub>) and the grain size of nanocrystals obtained using various methods [849]. Bulk-SPD [1275–1277] and physical vapor deposition (PVD) [1278, 1279] resulted in a significantly large improvement, as high as 0.8 V, in UFG materials although there were some exceptions [1280–1282]. Although surface-SPD methods such as SMAT can reduce the grain size to levels comparable to that of PVD [1267, 1283–1286], its effect on corrosion resistance appears to be less, probably owing to the presence of surface defects such as contamination or roughness. In Fig. 59b, the closed dots indicate that the chromium enrichment in the passive film after passivation was confirmed by X-ray photoelectron spectroscopy (XPS) or glow discharge optical emission spectroscopy (GD-OES) whereas the open dots indicate that chromium enrichment was not observed or confirmed. Despite a relatively large grain size and insufficient grain boundary area in the materials processed by bulk-SPD, uniform surface chromium enrichment might be achieved. The residual dislocations and their internal stress are the possible promoters of more uniform protective films over the entire surface. Fig. 59c shows the pitting potential of Fe-12Cr (wt%) alloys after eight ECAP passes and subsequent annealing [1275]. Note that the

pitting potential decreased appreciably after flash annealing at 698 K despite of negligible grain growth. The reduction of the peak-broadening in XRD after this flash annealing indicates that strain associated with residual dislocations, and possibly non-equilibrium grain boundaries was reduced. Thus, stress relaxation associated with structural recovery could degrade corrosion resistance.

The enhancement of the Cr/Fe ratio in the passivation film by the elastic or plastic deformation has been reported in Fe-Cr alloys [1287–1289]. One can consider the effect of dislocations on the Cr/Fe ratio based on the kinetics of the anodic dissolution of iron and chromium. For the anodic dissolution of the Fe-Cr solid-solution system, the degree of preferential dissolution of iron could be estimated by the ratio of the dissolution currents of iron and chromium. Under the stress field of dislocations, the interaction energy, w, arises between the solute chromium and the stress field and affects the dissolution currents of chromium and iron atoms differently according to the following equations.

$$i_{Fe} = A x_{Fe} \exp \left\{ - \frac{\Delta Q_{o,Fe} - (1 - \beta) F \mu}{kT} \right\} \quad (29)$$

$$i_{Cr} = A x_{Cr} \exp \left\{ - \frac{\Delta Q_{o,Cr} - (1 - \beta) F \mu - w}{kT} \right\} \quad (30)$$

where $x_{Fe}$ and $x_{Cr}$ are the atomic fractions of iron and chromium, respectively, $\Delta Q_{o,Fe}$ and $\Delta Q_{o,Cr}$ are the activation energy at equilibrium, F is the Faraday constant, $\beta$ is a symmetrical factor, $\mu$ is the potential drop, A is constant, k is Boltzmann's constant and T is temperature. A simple estimation of the interaction energy arising from atomic mismatch and modulus misfit between the solute chromium in iron indicates that the interaction energy becomes –0.06 eV at the core of the edge dislocations and –0.07 eV at screw dislocations [849]. The negative value of the interaction energy means that the solute chromium is stabilized near dislocations, which leads to lower dissolution kinetics than iron. Assuming that all solute chromium atoms are positioned at the core of screw dislocations, the selective dissolution, $i_{Fe}/i_{Cr}$, could reach $2.2 \times 10^4$, where $i_{Fe}$ and $i_{Cr}$ are anodic currents corresponding to the dissolution of iron and chromium atoms, respectively. This value is more than $10^2$ higher than that for no residual dislocations. The high-density residual dislocations of the order of $10^{15}-10^{16} m^{-2}$ in UFG can affect the dissolution balance between iron and chromium through their stress field and may promote the local chromium enrichment and protective ability of passive films [849].

### 6.10.2. Corrosion resistance of steels

As discussed in an earlier chapter, processing steels using SPD techniques enables obtaining UFG-structured steels with improved mechanical properties and corrosion resistance [1290, 1291]. The main mechanism for enhancing corrosion resistance by SPD is the formation of a passive protective layer that is more stable than in coarse-grained (CG) steels, mainly because of the high grain boundary density. The native oxide layer is formed in both coarse-grained and UFG steels, as shown in Fig. 60a. However, the high density of grain boundaries in UFG steels produces a uniform, compacted and well-adhered passive layer. This in turn results in uniform corrosion instead of localized corrosion because there are no preferential sites for cathodic or anodic corrosion [1276, 1292]. In addition, there are high-energy boundaries that also contribute to corrosion enhancement. Although corrosion in steels prefers to take place in the high-energy grain boundary region, there is also high atomic activity through diffusion to the surface as well as long diffusion paths. This influences the re-passivation process forming a thicker and more efficient passive layer [1293]. This mechanism has been also observed for other severely deformed metals [1258, 1294].

Different studies have been performed using SPD techniques to study the enhanced corrosion resistance over a range of steels. For carbon, alloyed and stainless steels, which have been mainly processed using
<figure><img src="image_138.png" /><figcaption>**Fig. 60.** (a) Differences between coarse grain (CG) and ultrafine/nanograined (UFG/NC) in the formation of native oxide passivation layer in steels [1276]. (b) Comparison of corrosion potential and current density between different steels processed by SPD: IF steel by ECAP [1295], austenitic steel by ECAP [1296], 304 stainless steel by ECAP [1297], 316 L stainless steel by ECAP [1276, 1298] and 316 L stainless steel by HPT [1299, 1300].</figcaption></figure>

ECAP and HPT, the stability of the protective layer has been mainly characterized in neutral and alkaline environments by the corrosion potential ($E_{corr}$), the corrosion current density ($I_{corr}$) and the pitting potential ($E_{pit}$). The corrosion potential indicates the tendency of the steel to corrode, and therefore it should be low. The corrosion current characterizes the rate at which the steel corrodes in a given environment and it also should be low [1301]. The pitting potential indicates the passivity layer breakdown and therefore, it should be high. The enhanced corrosion resistance and the stability of the passive layer in different steels processed using ECAP have been reported in low- and medium-carbon steels [1302, 1303], extremely low-carbon steels [1277], austenitic alloyed steels [1296], mild steels [1304] and interstitial-free (IF) steels [1295]. For stainless steels, although they are commonly characterized by good corrosion resistance, most of the literature also confirms that they can be improved considerably using SPD processing. In particular, the improved corrosion for stainless steels has been reported for 304 SS [1297], 316 L SS [1298, 1305] and F138 SS [1306] after ECAP processing and also for 316 L SS after HPT processing [1299, 1300]. For comparison, Fig. 60b shows some results for $E_{corr}$ and $I_{corr}$ in these steels. It should be noted that surface-SPD techniques such as SMAT [5] were also applied to low-carbon steels [1307], AISI 304 [1308], AISI 316 [1309, 1310] and AISI 316 L [1311]. By SMAT, a decreased grain size can be obtained along with high surface roughness [1312]; however, depending on the surface roughness and defects produced, the passive layer stability would be affected. In this sense, SMAT could be used for providing short-term corrosion protection to steels [1310].

Although there is a notable enhancement in corrosion resistance of steels processed using surface-SPD and bulk-SPD techniques, it is necessary to consider other factors that also influence corrosion

resistance, such as surface roughness, high density of defects and strain-induced martensite. For example, it is known that martensitic steels have lower corrosion resistance, because martensite promotes the pitting sensitivity of the steel, forming or accumulating crack nucleation sites, and as a result, it decreases the stability of the passive film [1313]. In this sense, the strain-induced martensite formed during the SPD process, especially for austenitic stainless steels deteriorates the corrosion resistance [1314, 1315].

## 6.11. Biomaterials

Biomaterials are defined as materials capable of being introduced into a biological environment to elicit a desired physiological response, such as the promotion of tissue growth or repair. Conversely, they are also expected to not induce any adverse effects, such as inflammation, cytotoxicity, or thrombogenicity [1316–1318]. Biomaterials are intrinsically related to the concept of biocompatibility, which can be defined as the ability of a material to perform with an appropriate host response in a specific application [1319, 1320]. It is important to emphasize that biocompatibility extends beyond the material itself; rather, it encompasses the entire material-host interplay [1321]. The properties of biomaterials are strongly influenced by their composition and grain size and they are available in various forms, including polymers, ceramics, metals and composites. Among these, metallic biomaterials stand out as a pivotal category, finding utility in a broad range of medical devices ranging from dental and orthopedic implants, heart valves and intraocular lenses to artificial hearts, biosensors and pacemakers [1322–1324]. Unquestionably, titanium and its alloys represented the most prominent group within the metallic biomaterials, mainly because of the desired combination of properties required for a biomaterial, such as high strength, low Young's modulus, high corrosion resistance, good fatigue properties, lightweight and good biocompatibility [1325–1328]. While titanium is of high interest for permanent implants, magnesium alloys are receiving significant attention for biodegradable implants. Despite the remarkable achievements with the utilization of titanium, magnesium and their alloys in their microcrystalline form in many different biomedical devices and components, the introduction of these biomaterials in the UFG form is considered the next smart generation of biomaterials [132, 1329, 1330].

When metallic biomaterials are scaled down to the nanometer level, their compatibility and bioactivity with osteoblast and fibroblast cells undergo a remarkable enhancement. This leads to a substantial increase in cell adhesion and proliferation on the surfaces of these materials, as evidenced by various studies [1331–1333]. Moreover, nanostructured materials exhibit higher levels of protein adsorption, cell attachment, proliferation and differentiation when compared to conventional materials [1333]. These physiological processes, involving the interaction of human cells with the surface of biomaterials, are notably enhanced at the nanometer scale, ensuring robust osseointegration [1331–1335]. It has been shown that a surface with nanoscale grain size exhibits higher bone formation efficiency compared with micro-grained surfaces [1334, 1335]. Within this context, SPD techniques have proven to be an outstanding and straightforward method for producing UFG biomaterials [1325, 1329, 1330, 1336]. Numerous studies have highlighted the enhanced properties of nanostructured titanium and its alloys when subjected to SPD methods compared to non-processed titanium. These improvements encompass superior corrosion resistance [1337, 1338], enhanced mechanical attributes including hardness, strength and wear resistance [657, 1329, 1339], enhanced bacterial adhesion [1340, 1341], improved wettability [1342, 1343] and high bioactivity and cell adhesion [1344–1346]. Notably, osteoblast cell proliferation on nanostructured titanium has been reported to be up to 19 times greater than on conventional titanium [1345]. Nanostructured titanium alloys, including Ti-6Al-4V [1347, 1348], Ti-6Al-7Nb [1349, 1350], Ti-Nb [312, 1351], Ti-Mo [1352, 1353] and Ti-Nb-Ta-Zr (TNTZ) [1354, 1355], have demonstrated remarkable mechanical properties and
biocompatibility enhancements when prepared using SPD methods such as ECAP or HPT. These materials have showcased advantages over microcrystalline forms of pure titanium and conventional titanium alloys [312, 1347–1355].

In the following sections, the effect of SPD on Ti- and Mg-based biomaterials is reviewed. Moreover, the contribution of SPD to the development of innovative biomaterials such as HEAs [317, 1356] and metal-protein nanocomposites [1357] is briefly discussed.

### 6.11.1. Biocompatible titanium alloys

Titanium and its alloys are known for their lightweight nature and excellent mechanical properties, corrosion resistance and biocompatibility. Compared to other major metallic biomaterials used in bone replacement implants, such as stainless steel and Co-Cr alloys, titanium demonstrates higher biocompatibility, making it a popular choice in the field of metallic biomaterials. Moreover, titanium stands out as the only metal capable of osseointegration, forming a bond with bone at the optical microscope level. Due to this unique characteristic feature, it finds extensive application in devices that interact with bone, including artificial hip joints, spinal fixation and dental implants. Presently, the most commonly used titanium alloy for biomedical purposes is the α+β-type Ti-6Al-4 V ELI (extra low interstitial), which boasts greater toughness than the Ti-6Al-4 V alloy [1358]. Pure titanium, such as commercially pure titanium grade 2 [538, 1359–1361] and grade 4 [1362–1364], which have lower strength than Ti-based alloys, have better biocompatibility than alloys, but they suffer from relatively low strength. There have been attempts to strengthen these materials by SPD which resulted in their commercialization. Here, some findings about SPD processing of Ti-based biomaterials are reviewed.

It was reported that the fatigue properties increased to 220 MPa from 180 MPa when etching the surface of UFG titanium grade 4 obtained through ECAP [1362]. Ti-based alloys exhibit high heat treatability, and their initial microstructure can be transformed into a lamellar structure or martensite through SPD, promoting grain refinement. Raab *et al.* conducted a study on titanium grade 4, characterized by a high impurity concentration among pure titanium materials, by quenching from the β region to form martensite [1363]. Subsequently, they subjected the materials to ECAP processing which resulted in a remarkable tensile strength of 1100 MPa and an elongation of 7.2%. The ECAP processing in the state of martensite further contributed to grain refinement by reducing the ability for grain boundary mobility, propagation of dislocations in grain interiors, and stress localization at interphase boundaries.

The α+β-type Ti-based alloy, such as Ti-6Al-4 V (wt%) alloy [1365–1371], Ti-13 V-3Al (at%) [1372], Ti-6Al-7 Nb (wt%) [1373, 1374], have been also studied by SPD. The strength can be increased in the α+β type Ti-based alloy by controlling the initial microstructure. In the Ti-6Al-4 V ELI alloy, Alagić *et al.* reported that both high tensile strength of 1546 MPa and total elongation of 18.8% can be achieved with HPT at 500 °C using the initial microstructure of approximately 90% α' martensite [1375]. In the Ti-6Al-7 Nb alloy, which was developed for medical use by replacing toxic vanadium with biocompatible niobium, Polyakova *et al.* reported 1460 MPa tensile strength and 11% of ductility after ECAP [1373]. Reportedly, 1280 MPa tensile strength and 22% of total elongation were obtained in Ti-6Al-7 Nb after HPT with the bi-modal initial microstructure of equiaxed α and α' grains [1374]. Since high Young's modulus leads to bone resorption, β-type Ti-based alloys with low Young's modulus, such as Ti-3 Nb [1376], Ti-10M-o-8 Nb-6Zr [1377, 1378], Ti-18Zr-15 Nb [1379, 1380], Ti-15Mo [1353], Ti-13 Nb-13Zr [1381, 1382], Ti-29 Nb-13Ta-4.6Zr (TNTZ) [1383], Ti-3.5Al-5Mo-4 V [1384] (all in wt%), have been processed by ECAP or HPT. Recently, Ti-containing HEA TiNbZrTaHf was processed by HPT which showed a high hardness of 564 HV with a low Young's modulus of 79 GPa and good biocompatibility of high entropy alloy after HPT [317].

Grain refinement not only enhances the mechanical properties of

<figure><img src="image_139.png" /><figcaption>**Fig. 61.** Histogram of target mRNA expression levels after 7 days of induction in MC3T3-E1 cells cultured on Ti-6Al-7 Nb (wt%) alloy before HPT (HPT-0), after HPT under 2 GPa (HPT-2) and after HPT under 6 GPa (HPT-6) in comparison with tissue-culture-treated polystyrene dishes (TCPS). Expression of (a) Runx2, (b) Col1α1 and (c) Akp2 was detected by real-time RT-PCR [1385].</figcaption></figure>

titanium but also improves corrosion resistance and biocompatibility. The corrosion resistance of UFG pure titanium grade 2 is enhanced through grain refinement by ECAP, leading to the formation of a passive layer on the titanium surface more effectively [1360]. Chen *et al.* reported that grain refinement of Ti-6Al-7 Nb alloy by HPT leads to enhanced osteogenic differentiation, as shown in Fig. 61 [1385]. The SPD process increases the fraction of grain boundaries and dislocations and positively changes the thickness and chemical component of the surface layer for biocompatibility [560, 1385]. All these findings suggest
that SPD techniques provide a powerful means for fabricating advanced biomaterials, holding immense potential for further innovation and development in the biomedical industry.

### 6.11.2. Biocompatible magnesium alloys

New-generation biodegradable implants possess the potential to supplant permanent medical implants as they assist fractured bones throughout the healing process and subsequently undergo controlled absorption by the body [1386, 1387]. Among the biodegradable metallic materials, magnesium alloys are the most relevant. The fundamental concept of designing Mg-based alloys as biomaterials begins with the formulation of alloy composition, followed by the implementation of manufacturing and processing techniques (e.g. utilizing SPD techniques) as well as surface modifications to achieve the desired

properties. This process culminates in the production of finalized products intended for long-term clinical trials, as depicted in Fig. 6a. Magnesium alloys with UFG morphology, produced via SPD techniques such as ECAP and HPT [5, 23, 32, 299, 1388], are especially interesting biomaterials for temporary orthopedic implants as they exhibit enhanced mechanical properties, adequate corrosion resistance and biocompatibility [1389].

ECAP processing of magnesium alloys is especially interesting as this technique can provide large-scale samples allowing the fabrication of bone screws and plates. When doing ECAP of magnesium alloys, not only grain refinement by shear but also dynamic recrystallization plays a substantial role due to the temperature rise during the process, stored energy in deformation-induced defects, and the low deformation rate [1390]. The extent of refinement depends also on the initial grain size.

<figure><img src="image_140.png" /><figcaption>**Fig. 62.** (a) Concept of designing Mg-based alloys as biomaterials. (b,c) Tensile strength and elongation at fracture for Mg-Zn-Ca alloys processed by (b) ECAP and (c) HPT (the numbers in parentheses for HPT-processed samples indicate the number of rotations and ‘HT’ indicates a subsequent heat treatment) [1389].</figcaption></figure>
Texture development during ECAP can lead to enhanced ductility and
softening effects [1391–1393]. Additionally, ECAP induces fragmenta-
tion and redistribution of hard particles within a magnesium matrix,
contributing to both, grain refinement and precipitation strengthening
[1394–1396]. The popular Mg-Al-based alloys like AZ31 [1397, 1398],
AZ91 [1399] or LAE442 [1400–1403] processed by ECAP show signif-
icant grain refinement and increased yield and tensile strength. How-
ever, the potential use of alloys with aluminum addition as
biodegradable implants is limited because the neurotoxic effect of
aluminum can cause Alzheimer’s disease or dementia when exceeding
the daily dose [1404, 1405]. Magnesium alloys containing zinc and
calcium are therefore in the research focus of different groups
[1406–1412]. They are characterized by excellent biocompatibility and
demonstrate attractive properties desirable for medical applications
without possible side effects [1406, 1408]. These alloys also exhibit
enhanced mechanical properties after ECAP, attributed to grain refine-
ment and precipitation hardening, as shown in Fig. 6b. Rare-earth el-
ements, in general, improve the mechanical properties of magnesium
alloys [1395, 1396, 1413–1415] and their corrosion resistance [1416],
with some alloys demonstrating good biocompatibility, such as the
WE43 alloy containing rare-earth elements and zirconium. However,
their content must be controlled to avoid the risk of harmful effects on
the human body [1417–1420]. The Mg-Ag alloys, known for their
antibacterial properties, exhibit increased strength and ductility after
ECAP processing [1421–1425], making an interesting option for
biodegradable medical implants to reduce the symptoms of local
inflammation or infection that can arise in the first weeks after
implantation.

HPT processing of magnesium alloys results in an initially inhomo-
geneous microstructure, but a gradual evolution towards relatively
uniform microstructures due to saturation tendencies occurs [177, 1138,
1426, 1427]. The number of turns required to achieve a homogeneous
microstructure depends on the applied pressure during HPT [1426,
1428]. Additionally, HPT induces a strong basal texture with the c-axis
parallel to the normal direction of the disc [1427, 1429, 1430]. Particles
in the magnesium alloy matrix undergo fragmentation and redistribu-
tion during HPT [1431], although the high hydrostatic pressure of the
process limits their breakage and spreading. Compared to ECAP, the
HPT process of magnesium alloys can be performed at lower tempera-
tures – even at room temperature – and the induced strain is generally
higher. For Mg-Al alloys, HPT processing results in significantly smaller
grain sizes (100–200 nm) compared to ECAP, thereby resulting in even
higher hardness values [1426, 1428, 1432–1434]. HPT-processed
Mg-Zn, Mg-Ca and Mg-Zn-Ca alloys (Fig. 62c), which are exclusively
alloyed with essential elements found in the human body, exhibit
improved mechanical properties through grain refinement and precipi-
tation strengthening [1409, 1435–1437]. Moreover, magnesium alloys
containing rare-earth elements [1438–1440], such as the WE43 alloy
[1441], exhibit increased hardness, yield stress and tensile strength after
HPT processing.

Magnesium alloys have high corrosion rates due to their non-noble
nature, which poses challenges for their use in biodegradable im-
plants. The corrosion mechanism involves the formation of hydrogen gas
and magnesium hydroxide and subsequent transformations, influenced
by factors like chlorine ions [1441] and the presence of impurities
[1409, 1442]. Alloying elements and thermomechanical treatments
affect corrosion and biodegradation, with galvanic coupling and
micro-galvanic corrosion playing a role [1443]. The influence of SPD
methods on biodegradation varies depending on the specific alloy and in
vitro testing conditions [1412, 1423, 1444–1452]. An increase in the
density of grain boundaries and especially changes in particle distribu-
tion after SPD processing can positively affect the corrosion rate of
magnesium alloys. *In vivo* studies on SPD-processed magnesium alloys
are limited but suggest improved degradation resistance and bioactivity
compared to untreated alloys [1453–1455].

In conclusion, magnesium alloys, which exhibit good

biocompatibility and homogenous and slow biodegradation, can
considerably be improved in terms of mechanical behavior by SPD
processing to be used for load-bearing medical implants in orthopedics.
A heat treatment or a coating to regulate the degradation rate in the first
stages of implementation might help to further improve their applica-
bility for biodegradable medical implants. Future in vivo studies are
needed to clarify the effect of SPD on the biocompatibility of these
materials.

6.1.1.3. *Innovative biomaterials*

The outstanding performance of nanostructured titanium and mag-
nesium in the biomedical field has inspired ongoing efforts to develop
novel biomaterials with even superior mechanical properties and
exceptional biocompatibility. Several studies have contributed to this
endeavor by harnessing the synergistic benefits of novel compositions
and nanostructuring achieved through SPD techniques. Notably,
emerging materials like HEAs [317, 1356] and Ti-protein composites
[1357] - especially designed for biomedical applications - have emerged
as promising candidates. These new biomaterial compositions processed
by the HPT method have presented mechanical properties and
biocompatibility fairly superior in comparison to traditional bio-
materials such as pure titanium, Ti-6Al-4 V, Ti-6Al-7 Nb, Co-Cr-Mo and
316 L alloys. As an example of this innovative approach, exceptional
properties including ultrahigh hardness, a low elastic modulus and
remarkable cellular activity for functions such as proliferation, viability
and cytotoxicity were reported for TiAlFeCoNi alloy after processing by
HPT [1356]. As shown in Fig. 63a, the average hardness witnessed a
notable increase from 635 Hv to 880 Hv by HPT processing due to grain
refinement. Additionally, the elastic modulus is 123 GPa which is
comparable to titanium, as shown in Fig. 63b. The cellular metabolic
activity of the TiAlFeCoNi alloy surpasses that of pure titanium and the
Ti-6Al-7 Nb alloy by an impressive margin ranging from 260% to
1020%.

Since SPD is capable of cold consolidation of powders, a recent study
processed the mixture of titanium and small amounts of an endogenous
protein BSA (bovine serum albumin) by the HPT method [1357]. This
new family of metal-protein nanocomposites exhibited an exceptional
combination of high strength and superior biocompatibility. The scan-
ning electron microscope (SEM) images, depicted in Figs. 64a and 64b,
show a reasonable mixing between the titanium and BSA particles at the
micrometer level and the elemental mapping of Fig. 64c confirms that
the black particles correspond to the BSA protein containing carbon
atoms which are surrounded by a titanium matrix. After HPT processing,
the BSA protein is mixed with titanium particles forming well-defined
protein layers that are distributed all over the sample, while the tita-
nium phase keeps its three-dimensional network, as displayed in
Fig. 64d-f. The cellular viability after HPT processing was evaluated by
*in vitro* cell culture experiments employing the MC3T3-E1 cells, as
indicated in the MTT assay in Fig. 64g. As can be seen, the
HPT-processed composites present superior biocompatibility in com-
parison with the coarse-grained titanium reference and the sample with
5 vol% of BSA showing the best biocompatibility performance among
other counterparts.

In conclusion, the introduction of innovative biomaterials, such as HEAs and Ti-protein composites, represents a significant stride in the biomedical field by the contribution of SPD. These newly designed materials offer enhanced mechanical properties and biocompatibility compared to conventional counterparts and can be further tuned for particular biomaterial applications.

7. Ultrafine-grained materials processed by SPD

Although SPD processing was preliminarily considered a metal
processing technology [1, 2], its application is now extended to a wide
range of metallic/nonmetallic materials, inorganic/organic materials
and traditional/innovative materials. SPD processing of metals goes
<figure><img src="image_141.png" /><figcaption>**Fig. 63.** (a) Variation of Vickers microhardness against distance from disc center for ingot and sample processed by HPT. For comparison, hardness for TiAlFeCoNi after compression (N = 0) and N = 1 turn of HPT, as well as the hardness range for Ti-based biomaterials, were included. (b) Nanoindentation load against displacement for HPT-processed TiAlFeCoNi sample [1356].</figcaption></figure>

back at least to the Bronze Age and SPD processing of organic materials should have a very old history in food making; however, SPD processing of hard and brittle materials was realized in the 1930 s following the introduction of the HPT method [20, 22]. Although SPD processing is mainly applied to metals (magnesium, aluminum, titanium, iron, etc.) and their alloys [1, 2], there are appreciable attempts to design new metallic materials (including, composites, new conventional alloys, HEAs and metallic glasses) using the SPD methods for different functionalities, as discussed in previous chapters. The application of SPD to semiconductors [1456], ceramics [1457] and polymers [14] also experienced significant progress in recent years, although the number of groups actively working on these materials is still limited. In the following section, research activities on various SPD-processed UFG materials are reviewed.

## 7.1. Magnesium alloys

The inherent brittleness of magnesium sets difficulties for room temperature processing. Therefore, early attempts for processing through ECAP [1458, 1459] and ARB [1460] were carried out at high temperatures. Processing routes which included extrusion followed by ECAP [1461] and the application of a back pressure in ECAP [1462] were developed, enabling the fabrication of UFG magnesium, but the minimum grain sizes achieved were in the range of several hundreds of nanometers. True nanocrystalline magnesium, with grain sizes of ~100 nm or smaller, is only produced via HPT processing at room temperature. Thus, the increase in research groups using HPT in recent years unveiled many interesting aspects of magnesium properties, such as the relationship between flow stress and grain size.

The early reported experimental data were limited to relatively coarse-grained samples and they usually fitted the well-established Hall-

<figure><img src="image_142.png" /><figcaption>**Fig. 64.** (a,b) SEM images in backscatter electron mode and (c) corresponding energy-dispersive X-ray spectroscopy elemental mappings of Ti + 5 vol% BSA protein before HPT processing. (d-f) SEM images at different magnifications in backscatter electron mode for Ti + 5 vol% of BSA composite after 5 turns of HPT under 2 GPa. (g) MTT cell viability assay examined by light absorbance at 570 nm for pure titanium and for nanocomposites containing 2 and 5 vol% of BSA protein produced by 5 turns of HPT under 2 GPa in comparison with hardness of coarse-grained annealed titanium [1357].</figcaption></figure>

Petch relationship in which the flow stress is inversely proportional to the square root of the grain size. The proportionality factor, Hall-Petch slope K, was usually high and exhibited a pronounced dependency on temperature [1463, 1464]. Recent studies show that this relationship is not constant for all grain sizes. A change in slope in plots of flow stress vs. grain size and grain refinement softening at small grain sizes have been reported [1465–1469]. A recent review [1470] evaluated multiple magnesium alloys and showed that a change in slope takes place at grain sizes of ~ 4 µm. The data is shown in Fig. 65 and two distinct regimes are observed. Hence, data for coarse-grained magnesium alloys follow a larger slope attributed to twinning-controlled deformation, while data for ultrafine and nanocrystalline follow a reduced slope associated with slip-controlled deformation. It is worth noting that the data in the range of slip-controlled deformation agree with the mechanism of grain
<figure><img src="image_143.png" /><figcaption>Fig. 65. Flow stress and elongation to failure of magnesium alloys plotted as a function of the grain size [1470].</figcaption></figure>

boundary sliding [1471] in which the flow stress is predicted by Equation 31. This mechanism has shown good agreement with experimental data for multiple materials with different grain sizes and tested at different temperatures and strain rates [739, 1472, 1473].

$$ \sigma \approx \sqrt{\frac{3GkT}{2d,b^2} \ln\left(\frac{\dot{\epsilon}d_s^3}{10\delta D_{gb}} + 1\right)} \quad (31) $$

where $\sigma$ is the flow stress, $G$ is the shear modulus, $k$ is the Boltzmann's constant, $T$ is the absolute temperature, $d_s$ is the spatial grain size, $\dot{\epsilon}$ is the strain rate, $b$ is the Burgers vector, $\delta$ is the grain boundary width and $D_{gb}$ is the coefficient for grain boundary diffusion. The data in Fig. 65 shows that high strength, in the range of 400 MPa, might be observed in UFG magnesium alloys. The highest strength in a magnesium alloy was reported in a Mg-17Ni-17 Pd (at%) alloy with the BCC structure processed by HPT [296] but this value is not shown in Fig. 65. A noticeable increase in elongation to failure is observed in the grain size range 2–8 µm which is the range for transition in deformation mechanism. Thus, a good combination of strength and ductility is observed in this grain size range.

Early papers showed that ECAP can introduce superplastic behavior in magnesium alloys [796, 798, 1474–1477]. In practice, grain refinement enables a decrease in temperature and/or an increase in strain rate for superplasticity. Also, the elongations reported in magnesium processed by ECAP tend to be larger than in conventional thermo-mechanical processing. Fig. 66a shows the appearance of tensile specimens of a ZK60 alloy processed by different numbers of passes of ECAP and pulled to failure at 473 K. Exceptional elongations of over 1000% were reported, including an elongation of 3050% [1475].

It was shown later that superplasticity is also observed in magnesium processed by HPT [801, 1479, 1480] and analyses of experimental data showed that the strain rate for superplasticity agrees with theoretical predictions from the mechanism of grain boundary sliding [349, 1481]. In addition to the development of high-temperature superplasticity, severe plastic deformation may also introduce exceptional ductility in magnesium at low temperatures and even at room temperature. For instance, superplastic elongations were reported in an AZ91 alloy tested at 423 K [1480] and in an Mg-Li alloy tested at 373 K [800]. Although many magnesium alloys display brittle behavior at room temperature, extraordinary ductility has been reported in UFG pure magnesium [1482] and the Mg-8Li (wt%) alloy processed by HPT [350]. Fig. 66b shows the stress-strain curves of the Mg-8Li (wt%) alloy tested in tension at room temperature in the extruded condition and after 5 and 200 turns of HPT. An increase in elongation is observed after 5 turns of HPT and further processing to 200 turns increases the elongation to over 400%, which is considered a superplastic elongation [350].

As discussed earlier, SPD processing also improves the hydrogen storage performance of magnesium [1077]. It was shown that ECAP processing increases hydrogen desorption kinetics [1143]. It was also demonstrated that HPT processing improves hydrogen storage capacity

<figure><img src="image_144.png" /></figure>

<figure><img src="image_145.png" /></figure>

<figure><img src="image_146.png" /><figcaption>Fig. 66. (a) Appearance of tensile specimens of a ZK60 alloy processed by different number of passes of ECAP and tested at 473 K [1475]. (b) Stress vs. strain curves and appearance of tensile specimens of an Mg-8Li (wt%) alloy processed by extrusion and different number of turns of HPT and tested at room temperature [350]. (c) The corrosion rate of magnesium and magnesium alloys plotted as a function of the flow stress [1478].</figcaption></figure>

[1138]. A significant advance in this field was achieved by mixing magnesium with other metals using ultra-SPD through many turns in HPT. Thus, a Mg₄NiPd alloy with BCC structure processed by 1500 turns of HPT exhibited room-temperature hydrogen storage capability [359]. The hydrogen pressure composition isotherm for tests at 305 K under hydrogen pressures up to 10 MPa shows up to 0.86 wt% of hydrogen adsorption after the fifth cycle with 0.7 wt% reversibility [359]. Besides the improvement in hydrogen storage, it has been suggested that grain refinement can induce pseudo-uniform corrosion due to the increase in the volume fraction of grain boundary area [1258]. Experiments in pure magnesium [1483] and a magnesium alloy [1484] processed by HPT support this suggestion. Also, many papers [1485–1488] reported a decrease in the corrosion rate of magnesium and its alloys after SPD.
These observations suggest that SPD processing can improve the performance of magnesium for biological applications as temporary implants. Fig. 66c shows experimental data of flow stress and corrosion rate for magnesium before and after SPD processing, indicating that HPT processing can increase the strength and decrease the corrosion rate of magnesium [1478].

## 7.2. Aluminum alloys

Since the early works of the late 20th century and early 21st century on the processing of UFG materials by SPD techniques, much attention has been paid to Al-based materials. They possess a highly attractive combination of functional properties: outstanding strength-to-weight

ratio, fair corrosion resistance, notable formability and so on. In addition, aluminum belongs to a limited number of elements that might be used as conductors being the cheapest variant among all. Therefore, Al-based alloys are widely used as constructional and conductive materials in transportation, food, construction, sport and other industries. However, their strength is relatively low while advanced applications require Al-based alloys with improved mechanical and functional performance. Thus, the enhancement of their mechanical and functional properties through nanostructuring by SPD is an important task to develop novel high-performance Al-based alloys for innovative applications. Fundamental studies and recent achievements in this field have been analyzed in the most recent reviews [835, 1489] and are summarized below.

The most important motivation for the application of different SPD

<figure><img src="image_147.png" /><figcaption>**Fig. 67.** (a) A possible modification of the Hall-Petch relationship for UFG FCC metals and solid solutions [1497]. (b) Evidence on grain boundary sliding at room temperature in UFG pure Al, shown by three-dimensional atomic force microscopy (AFM) [1498]. (c) The Al/Al boundaries strongly segregated by solute zinc atoms, causing wetted grain boundaries in SPD-processed Al-30Zn (at%) alloys [351]. (d) High-resolution TEM image showing a wetting layer along an Al/Al grain boundary in the HPT-processed Al-30Zn (at%) sample [1499], (e) Surface morphologies of the compressed pillars on the surface of the HPT-processed Al-30Zn (wt%) sample, demonstrating a direct evidence for the occurrence of intensive grain boundary sliding without any deformation localization [1500].</figcaption></figure>
methods comes from the 70-year-old Hall-Petch relationship [661, 1490], describing the improvement of the yield stress with decreasing grain size. The possibility of achieving high-strength states following the Hall-Petch equation, engineering superplastic behavior and other properties of metallic materials by forming UFG structures has traditionally been the motivation for research and development in the SPD field [1491]. Consequently, several early efforts have been made to refine the grain size in aluminum [1492–1495] and its alloys [348, 810, 817, 820, 1492]. Presently, the use of various SPD techniques has become the main approach in obtaining different UFG materials with nanoscale and submicron grain sizes. An important factor for UFG materials is that depending on the SPD processing regimes, different types of grain boundaries are formed in metals and alloys, including aluminum materials, such as low- and high-angle grain boundaries, equilibrium and non-equilibrium ones, as well as the grain boundaries containing significant segregations of alloying elements, which typically has a very strong influence on their properties [304, 373, 1496]. Besides, experimental experience indicates that the conventional Hall-Petch equation cannot be typically maintained in its original form for UFG materials. A possible modification of the original Hall-Petch equation has been suggested [1497] for a uniform description of submicron-sized FCC metals and solid solutions (Fig. 67a).

Together with the strengthening effect, the average grain size also plays an important role in the plastic deformation process, such as superplastic flow, where the main deformation mechanism is grain boundary sliding [791, 1501–1503]. Experimental evidence [1498] on room temperature grain boundary sliding in UFG pure aluminum deformed by Vickers indentation has been detected by using three-dimensional atomic force microscopy (Fig. 67b). Analysis of the pileups formed around the Vickers pattern suggests that the contribution of grain boundary sliding to the total strain in the UFG aluminum sample can reach up to 70% [1504]. Besides the grain-refinement, SPD may result in strong decomposition of the microstructure [351, 817, 1500, 1505–1507] and intensive segregation of solute atoms to grain boundaries (Fig. 67c) [351, 1499, 1508], causing so-called wetted grain boundaries [351, 1499, 1507–1509] in UFG alloys (Fig. 67d). It has been shown that segregated grain boundaries have an enhancing effect on grain boundary diffusion, facilitating grain boundary sliding, as shown in Fig. 67e, and leading to super-ductility [817] or even superplasticity [351, 1510] at room temperature in Al-Zn alloys. The grain boundary segregation also has a unique effect [812] in the case of UFG Al-Zn-Mg-Zr (7xxx) commercial alloys, where conventional materials can show superplastic behavior only at a relatively high homologous temperature of about 0.7–0.8 $T_m$ [315, 810, 814, 1511–1513]. Recent results have demonstrated the significance of enhanced diffusion along grain boundaries at low temperatures in a UFG Al-Zn-Mg-Zr commercial alloy, leading to its superplasticity with a record elongation higher than 500% at 443 K (0.47 $T_m$) (Fig. 68a), and maintaining relatively high strength after superplastic deformation [1510]. Considering the temperature dependence of the superplasticity of commercial aluminum alloys, the summary graph in Fig. 68b shows the significance of the mentioned new results obtained on the ultralow-temperature superplasticity of the UFG Al-Zn-Mg-Zr alloy.

Combining high mechanical strength with high electrical conductivity has always been a challenge for the aluminum industry, especially for the design of power transmission lines [1514] or for the development of advanced wirings for automotive applications [1515]. Alloying elements in solid solution or precipitates increase the strength but significantly scatter electrons giving rise to an increase of the electrical resistivity. The best compromises achievable by traditional treatments have been demonstrated with 6xxx alloys [857] but the design of UFG structures by SPD processes opens principally new possibilities to expand the limits of the trade-off between strength and conductivity in conductor alloys [834, 835]. There are however some limitations: (i) UFG structures require some stabilizers, solutes or particles that pin grain boundaries and limit grain growth [310]; (ii) grain boundaries are

<figure><img src="image_148.png" /></figure>

<figure><img src="image_149.png" /><figcaption>**Fig. 68.** (a) Deformation curves of an UFG Al-Zn-Mg-Zr (7xxx) alloy, showing a record total elongation higher than 500% at 443 K (homologous temperature: 0.47) [812]. (b) Significance of the UFG structure having grain boundaries segregated by solute atoms, resulting in ultralow-temperature superplasticity in the Al-Zn-Mg-Zr alloy (denoted as Present Study) [812].</figcaption></figure>

known to scatter conduction electrons and when the grain size is below the free mean path of the electrons, the electrical resistivity is significantly affected [1516]. Different strategies involving SPD have been proposed to overcome these barriers. The most straightforward approach (Fig. 69a) is to combine precipitation hardening and grain boundary strengthening in traditional conductor aluminum alloys (such as Al-Mg-Si [872, 1517–1520] or thermally stable Al-Zr-based alloys [884, 1521, 1522]). SPD might be carried out in the solutionized state or after precipitation treatment. In the first case, however, under SPD conditions the crystalline defects and strain-induced segregations [1519] may affect the precipitation and specific heat treatments and/or doping with other elements are required [872, 1518, 1521, 1522].

Alternatively, the potentiality of SPD to tailor new innovative systems has been explored, such as immiscible Al-RE (RE: rare earth) [315, 881, 1523], Al-Fe [310, 869, 877, 879, 880, 1524–1526] or Al-Ca [311, 1527] alloys. As-cast structures with a high-volume fraction of fine intermetallic particles or solute segregations (Fig. 69b) promote the grain refinement and the stability of UFG structures with limited effect on the electrical conductivity. Besides, SPD can additionally expand the limits for the structural design of Al-based alloys by forming grain boundary segregations providing extra-strengthening of the UFG state [1528] or by creating supersaturated solid solutions of elements that are immiscible under near equilibrium conditions [311, 315, 877, 881,
<figure><img src="image_150.png" /><figcaption>**Fig. 69.** (a) A strategy to improve the strength-conductivity ratio in age-hardenable aluminum alloys by SPD-controlled nano-precipitation [835]. (b) Grain boundary segregations in SPD-processed immiscible Al-RE (RE: rare earth) alloys [1523]. (c) A theory-based predictive approach to define limits of the hardness-conductivity trade-off in UFG Al-Mg-Si alloys as a function of the grain size and with 0.2 wt% of solutes left in solid solution. Different curves correspond to different contributions considered for the estimations: without precipitates (blue), with precipitates with density corresponding to the conventional T6 (red), with precipitates and maximal measured dislocation density after HPT at room temperature (orange). Green and orange squares correspond to experimental data collected on the 6101 and 6201 alloys, respectively [1519]. (d) A potentiality of SPD to expand the space of improving the strength and conductivity of conductor aluminum alloys [835]. IACS stands for International Annealed Copper Standard [857].</figcaption></figure>

[1524], which can transform these alloys to a type of age-hardenable ones. Basic theoretical estimations accounting for structural contributions to yield stress and resistivity can predict the limits of the strength-conductivity ratio for SPD-engineered Al-based alloys as shown in Fig. 69c [1519], but these estimations need to be modernized considering the latest findings.

Note that the morphology of the intermetallic phase in the as-cast state can severely affect the SPD efficiency for immiscible systems [880, 1523]. Here a combination of advanced casting techniques (such as electromagnetic casting [1529]) to disperse intermetallic eutectics before SPD can be fruitful [879, 1523, 1526]. As a result, electromagnetically cast bulk specimens of Al-Fe alloys after complex SPD-assisted treatment demonstrate enhanced thermal stability similar to Al-Zr conductors with a strength/conductivity ratio comparable to advanced Al-Mg-Si alloys [879, 1526]. Moreover, metal matrix composites processed by continuous SPD methods also offer some unique possibilities to achieve conducting wires [1530–1534]. Composites combining different aluminum alloys or Al-Cu alloys are more complex to produce, but extend the possibilities of strength/conductivity combinations. The full range of property combinations that could be expanded using the

different SPD-based approaches is summarized in Fig. 69d (based on the literature data collected in [835]). It is worth noting that there are several issues in the SPD-driven development of high-strength lightweight conductors, which are still poorly documented in the literature. Limited studies address such important properties as fatigue [876, 1515], creep [737, 1535] and corrosion [1515, 1536] behavior of electrotechnical UFG aluminum alloys, vitally important for their industrial applications. Despite the fewer promising results, these areas still wait for comprehensive exploration.

Results obtained on SPD-processed UFG aluminum alloys clearly show the innovation potential of unique grain boundary phenomena, such as the advantage of intensive grain boundary sliding at room temperature for micro-devices, formation of special grain boundary segregations for the combination of high strength and low-temperature superplasticity in commercial aluminum alloys. There is also obviously an open field for the design of unique UFG structures in lightweight Al-based alloys to achieve a full range of exceptional combinations of mechanical strength and electrical conductivity and multifunctional properties [180].
### 7.3. Titanium alloys

Titanium and its alloys have a special positioning in the SPD field because one of the most important successes in the commercialization of SPD is the use of Ti-based materials for biomedical applications. Application of SPD to pure titanium results in grain refinement to the submicrometer level and enhancement of strength over 1000 MPa. When pure titanium is processed by HPT under high pressure (usually over 4 GPa), it shows a phase transformation to the α phase which is hard and brittle. SPD is also widely applied to various titanium alloys with different initial phases, leading to grain refinement and phase transformations. Most studies on Ti-based alloys are directed toward biomedical applications by studying the mechanical properties, biocompatibility and corrosion resistance. This section reviews some recent findings on the application of SPD to titanium alloys with a focus on biomedical applications.

Prosthetic devices (Fig. 70a) are growing fast in commercial and medical importance, but progress in implant materials is limited [1537, 1538]. Developing new biomaterials for implants is challenging due to conflicting requirements [1327], as also described in earlier chapters. Prostheses must be biocompatible, free of toxic elements like vanadium, nickel, copper and cobalt [342, 1539, 1540], promote bioactivity, and have high corrosion and wear resistance. They should also possess superior mechanical properties, like a small Young's modulus similar to bone tissue to prevent stress shielding [1541–1544]. Titanium and its alloys are ideal for dental and orthopedic implants due to their biocompatibility and corrosion resistance. They possess a natural oxide

layer [1545] that protects against corrosion but does not promote bone healing or combat infections [1546]. Surface modifications and coatings can improve bone formation and reduce the risk of infections [556, 1350, 1547–1556]. Ti-based implants display exceptional mechanical properties (maximum load, bending, stiffness and fatigue strength) [1547, 1548, 1557]. These properties may effectively promote bone cell phenotypic needs [1558–1560]. Moreover, they have reasonably similar properties to bone tissue [1560]. Pure titanium is corrosion-resistant and is the best material so far for applications such as dentistry. Ti-6Al-4 V alloy has been chosen as a biomaterial because of its mechanical properties; however, there are concerns about its high elastic modulus [1561] and the cytotoxicity effects of aluminum and vanadium ions that can be released into the body due to corrosion or friction [342] and result in adverse bone tissue reactions [342, 1539, 1540]. Therefore, it is imperative to design single-phase β titanium alloys with high strength that are completely non-toxic and free from allergens (β alloys possess a lower elastic modulus intrinsically than α-Ti and the α+β Ti-6Al-4 V alloy) [342, 1557]. More comprehensive and up-to-date insights into the effect of thermomechanical processing on the microstructure and mechanical behavior of Ti-based alloys are given elsewhere [342, 1327].

Grain refining via SPD processes like ECAP and HPT can enhance the mechanical properties of Ti-based alloys. These methods effectively reduce grain sizes to the UFG or even nano-size range [1, 2, 177, 299, 556, 657, 668, 1364, 1549, 1550, 1562, 1564–1567], leading to improved strength, wear resistance and bioactivity. In addition to mechanical strength and stability, appropriate surface

<figure><img src="image_151.png" /><figcaption>**Fig. 70.** (a) Examples of prosthetic devices for joint and teeth replacement. (b) Average grain size and (c) hardness in titanium alloys Ti-6Al-7 Nb (TAN), Ti-13-Nb-13Zr (TNZ) and Ti-35 Nb-7Zr-5Ta (TNZT) processed by HPT under different pressures, where 2 t, 3 t and 5 t refers to the number of HPT turns and SC, FC and QA refer to the initial condition before HPT processing (samples were heated treated at 1023 K for 1 h and cooled down; SC: slow cooling from the β phase region in the furnace; FC: fast cooling from the β phase region by water quenching; QA: quenching, aging at 773 K for 5 h and cooling in air). (d) Young's modulus for undeformed and HPT-processed samples of Ti-6Al-7 Nb, Ti-13 Nb-13Zr and Ti-35 Nb-7Zr-5Ta alloys [556, 1562, 1563].</figcaption></figure>

<figure><img src="image_152.png" /></figure>
nano/micro-roughness is crucial for producing reliable implants with a low risk of failure incidence [1568]. The high strength of UFG implants allows for smaller sizes and less invasive surgeries while maintaining functionality [851]. In addition to grain refinement, when processing titanium and its alloys through SPD, phase transformations can occur. The formation of metastable phases such as the ω phase and martensite α" phase [556, 562, 1550, 1569–1571] can impact the tenacity and elastic modulus of titanium alloys [556, 1569, 1572, 1573]. Implants also require stress shielding and corrosion resistance and high bioactivity for a sufficient bone-bonding ability. This can be achieved by creating UFG microstructures through SPD and applying extra surface modifications such as anodization and acid + alkaline treatment.

Different titanium alloys like Ti-6Al-7 Nb (TAN) [1562, 1567], Ti-13-Nb-13Zr (TNZ) [556, 1550, 1566], Ti-35 Nb-7Zr-5Ta (TNZT) [556] and commercially pure titanium [1550, 1553] were subjected to SPD with varying strain rates, applied loads, starting phases and morphologies. An overview of the results of HPT processing, including phase transformations, structural/microstructural characterization, mechanical and corrosion properties, surface treatments (anodization and acid + alkaline treatment) and bioactivity has been compiled in [1563]. Deformation by HPT induces phase transformation in these alloys, which depends on the amounts of α or β stabilizers, the strain rate, applied load, starting phases and α phase morphologies [556, 1550, 1562, 1563, 1566]. When there are high levels of α stabilizers like oxygen or aluminum and α' present, a phase transformation to the ω phase is prevented. If α' is present, the fraction of the β phase increases with decreasing applied pressure. At high pressures, an α phase with a lamellar morphology transforms into the ω phase, while at low pressures, a fine globular α phase transforms into the ω phase (at higher

pressures, β can transform to α as well). The process of HPT is highly effective in achieving high grain refinement in Ti-based alloys down to grain sizes of 120 nm (Fig. 70b) leading to a maximum hardness of 1.4 GPa (Fig. 70c) [556, 1562, 1563]. A Hall-Petch analysis revealed that the strength of titanium alloys is not solely determined by the number of grain boundaries but also by specific characteristics, including different phases and their sizes [1563]. However, Young's modulus is solely affected by the phases, and an increase in the fraction of the β phase leads to a decrease in modulus (Fig. 70d).

To enhance the surface properties of SPD-processed titanium alloys, oxide nanotubes were made via anodization on α or α' phase surfaces [1563]. It was reported that the β phase produces oxide nanopores (thicker-walled tubular morphology), while treatment by acid + alkaline media can change the morphology to sponge-like structures [556, 1350, 1550]. The corrosion properties (Figs. 71a and 71b) indicate that the HPT-processed samples perform the best, although all cases have low current densities, indicating effective control of mass transport through the oxide layer [1550, 1553, 1554, 1563]. Corrosion potentials increased significantly with anodizing the surfaces, suggesting an effective strategy for the further improvement of SPD-processed alloys. Regarding bioactivity (Figs. 71c and 71d), polished surfaces are not bioactive due to their hydrophobic nature. In comparison to hydrophilic anodized alloys, commercially pure titanium hydrophilic anodized surfaces are more bioactive, regardless of the deformation of the surface. Morphologically, surfaces containing nanopores + nanotubes or only nanopores are not promising as well as pure titanium for apatite precipitation. Surface charges have a negative impact on precipitation for both anodized and chemically treated surfaces. For chemically treated surface tentacles to promote apatite formation, a critical dimension is

<figure><img src="image_153.png" /><figcaption>**Fig. 71.** (a,b) Corrosion behavior shown by potentiodynamic polarization curves recorded in simulated body fluid (SBF) at 310 K for (a) Ti-13-Nb-13Zr (TNZ) and (b) Ti-35 Nb-7Zr-5Ta (TNZT) alloys compared to commercially pure (CP) titanium with and without HPT processing. (c,d) Mass gain of apatite after immersing samples in SBF as a function of the soaking time, comparing treated and untreated surfaces of (c) TNZT and (d) TNZ before and after HPT processing [556, 1550, 1553, 1554, 1563].</figcaption></figure>
necessary for biological fluid bonding. Grain refinement is a crucial factor in surface bioactivation, as proven by the superior results attained by anodized and HCl-treated surfaces of HPT-deformed samples [556, 1550, 1563].

Nanostructuring by SPD effectively improves the properties of titanium and its alloys, particularly for biomedical applications [1563]. While HPT is limited in sample size for medical implant production [53, 177], alternatives like ECAP [1350, 1543, 1549] and ECAP-conform [156, 193] are better suited for commercial processing of Ti-based materials for biomedical applications [1330, 1574].

### 7.4. Iron alloys and steels

Iron and steel alloys are the most-used engineering metallic materials. These materials have also received high attention in the SPD field and several SPD techniques have been applied to them over the last decades. SPD processing of pure iron, its strength can exceed 1 GPa due to grain refinement. SPD has been applied to various iron alloys and steel. The application of SPD to austenitic stainless steel usually results in martensitic phase transformation and the reduction of grain size below 100 nm. Among iron alloys, carbon-based steels including pearlitic, bainitic and martensitic grades are of special interest to the SPD community because they play a crucial role in our society considering typical engineering applications such as bearings or gear wheels or products for the mobility sector (e.g. for railway transportation). In many of these applications contact fatigue and wear arise and may lead to great microstructural changes of the steels [1575]. As a consequence, defects may form and restrict the lifetime of components [1576–1578]. The microstructural changes can be experimentally simulated using SPD in a well-defined manner, as far as the applied strain and temperature are concerned [1579]. In this sense, it represents an industrial application of SPD for testing and ranking materials exposed to contact fatigue. The foundation for this SPD contribution is analyzing the microstructural evolution of such steels for the understanding of material strengthening and changes in ductility [1580], fracture resistance [1581] and fatigue crack growth behavior [1582] along with the connected mechanical anisotropy. This section provides an overview of the underlying structural and mechanical changes of pearlitic, bainitic and martensitic steels subjected to SPD, particularly to HPT [1583].

Although cold wire drawing is sometimes not considered an SPD technique due to the changing sample dimensions [5], it must be mentioned in the context of SPD of steels due to (i) its importance in

industry [1584], (ii) its close ability to deform materials of high hardening capacity [1585] and (iii) its status leading to the currently strongest pearlitic steels and therefore the strongest engineering alloy in the world [1586]. Cold-drawn pearlitic steels belong to the best-studied SPD materials and have been studied systematically with the pioneering work of Emburry and Fisher in 1966 [1587]. This production route provides thorough insights into the structural evolution up to the highest deformation strains. Different methodologies were used to reveal the loss of the initial colony structure, consisting of alternating ferrite and cementite lamellae, using X-ray-based techniques [1588], TEM [1589] or atom probe tomography [682], as presented in Fig. 72.

The transfer to a single-phase ferrite nanograin structure is accompanied by a severe change in ductility [682] and toughness [683]. Similar observations have been made by studying the structural evolution of pearlitic steels during HPT [1590], although the highest SPD strains are to date not accessible due to experimental restrictions [1591]. However, in any application field in which pearlitic steels and their components are exposed to severe strains (e.g. in rails [1592]), one has to deal with these structural changes to better forecast the lifetime or even actively counteract the deterioration of the lifetime by novel alloy design concepts. Deformation in a quasi-constrained setup by HPT also enables to study of the structural changes in bainitic [1593] and martensitic steels [1594]. Especially for the latter case, its inherent brittleness does not allow for cold drawing to elevated strains at ambient conditions. Yet, martensitic steels are imposed to severe straining in bearings, where so-called white etching layers limit their lifetime [1595]. Thus, studying their structural changes at high strains provides important insights into the basic deformation and failure mechanisms during service.

A recent overview demonstrated that severe deformation of the different initial steel microstructures leads finally to a very similar architecture as the limit of deformation, as summarized in Fig. 73 [1583]. The microstructures consist of elongated nano-sized ferritic grains decorated with carbon that stabilizes grain and phase boundaries [1596]. These findings show the importance of carbon distribution, as it has been investigated in martensite for decades [1597]. The evolution in the low and intermediate deformed regime is however strongly dependent on the initial structure, even if the ferrite composite structure consists of a percolating lamellar or randomly distributed carbide structure. It is astonishing that despite the lower carbon concentration in a martensitic Ck10 steel (~0.1 wt% carbon), only slightly coarser lamellar spacings can be achieved in ferrite, as compared to fully

<figure><img src="image_154.png" /><figcaption>**Fig. 72.** Evolution from an initially lamellar structure (strain: $ε_d$ = 1.96) to a ferritic subgrain structure with carbon either supersaturated in the matrix or segregated to subgrain boundaries (strain: $ε_d$ = 6.52) by straining [682].</figcaption></figure>
<figure><img src="image_155.png" /><figcaption>Fig. 73. Schematic illustration of the structural evolution of (a) pearlitic, (b) bainitic and (c) martensitic structures during quasi-constrained HPT deformed up to current technological limits [1583].</figcaption></figure>

pearlitic steel (0.6–0.8 wt% carbon), at the highest HPT strains. These similarities may indicate that white etching layers in pearlitic and martensitic steels, which are subjected to severe strains during service, might be related to similar mechanisms. Along with the microstructural similarities, the mechanical properties of pearlitic and martensitic steels show comparable tendencies within the deformation limits [1596, 1598], which highlights the importance of grain architecture for mechanical properties. Despite these clear trends, there are indications that the interface structure plays a significant role in plasticity, strength and fracture [1599–1601], and thus, further scientific scrutiny is required. The current major restriction to push forward the structural refinement of these steels is however the strength of the HPT anvil materials and the applicable hydrostatic pressure.

## 7.5. Lattice-softened alloys

Furuta *et al.* reported in 2015 that severely cold worked Fe-24.6Ni-5.8Al-0.4 C (wt%) had a yield strength of 2 GPa and a fracture elongation of 20%, [1602] which attracted attention since such a combination of strength and ductility outperforms the conventional trade-off. They designed the alloy composition by considering the phase stability, where they controlled the averaged valence electron number to generate lattice softening. With these mechanical properties, they also reported that the Lüders-type band propagates multiple times during tensile tests, resulting in high ductility. After the first report by Furuta *et al.*, several high-strength lattice-softened alloys such as Fe-Ni-Al-C, Fe-Mn, Fe-Cr-Ni, Fe-Ni-Mn have been reported to show similar Lüders-type deformation when they are processed by cold working [1603–1613]. Fig. 74a shows the relationship between tensile strength and total elongation, while Fig. 74b shows the relationship between stress of Lüders deformation, $\sigma_L$, and Lüders strain. Other reports have been made on these high-strength alloys [1614–1625], but among them, the Fe-Ni-Al-C alloy shows the best strength-ductility balance. There have been limited numbers of reports on the effects of alloy compositions and cold working conditions on mechanical properties and the propagation behavior of the Lüders bands in high-strength steels [1606–1608, 1610–1612, 1626]. However, it has been reported that the stability of the γ phase affects the magnitude of Lüders strain, while the strategies for controlling the microstructure to achieve high strength and high ductility are currently unknown.

Two publications recently reviewed the features of high-strength lattice-softened alloys [1628] and high-strength alloys with Lüders deformation [1627]. Recent experiments attempted to increase the strength of these lattice-softened alloys further by compositional modifications [1627]. A higher amount of aluminum increases the strength

and decreases the ductility in Fe-25Ni-(5.0–6.0)Al-0.3 C (wt%) alloys. The addition of aluminum increases the volume fraction of the B2 phase, Ni-rich second phase, which decreases the γ-phase stability of the matrix and accelerates transformation from γ to α' during cold rolling and tensile deformation. It was also confirmed that the γ-phase stability affects the Lüders strain. Fig. 74c shows tensile stress-strain curves of the specimens after cold rolling by 60% (a-n represent points where Lüders type deformation was confirmed). The alloy with 5.0 wt% Al has a yield point of about 1000 MPa and a total elongation of about 39%, and the alloy with 5.5 wt% Al, has higher strength and lower ductility. Increasing aluminum to 6.0 wt% leads to a higher deformation stress, but the sample fractured at a lower strain. The profiles for local strain along the center line parallel to the tensile axis in the gauge section were determined by digital image correlation (DIC) analysis at each point (a-f in the stress-strain curves of the alloy with 5.0 wt% Al and 60% cold rolling), as shown in Fig. 74d. The local strain increases near 10 mm of the gauge section at point a, after the yield point drop; the local strain profile changes as the nominal strain increases to point c; and the local strain increases and the deformation area expands over the entire test specimen, indicating that typical Lüders-type deformation proceeds. Similar Lüders-type deformation can be seen during the deformation at points d-f. In addition, the higher rolling reduction increased the Lüders strain. For example, severe cold rolling by 80% enabled the prolonged Lüders strain as much as 25% in nominal strain. This prolonged Lüders strain was achieved by multiple propagation of Lüders-type bands.

There exists specific phase stability of the γ phase and deformation texture that enables sustained and prolonged Lüders deformation in lattice-softened alloys. The importance of phase stability on deformation behavior has been also reported in other lattice-softened alloys [1629, 1630]. The stable Lüders deformation suppresses subsequent uniform deformation with large work hardening, which raises the deformation stress and promotes premature fracture. Although the Fe-Ni-Al-C alloy has interstitial carbon, the simple dislocation-locking mechanism cannot explain multiple propagation of Lüders-type bands. Besides the mechanism of dislocation locking, the Lüders deformation has also been found in UFG materials including pure metals [715, 1631–1635] and alloys [1636, 1637], which do not usually show the yield point phenomenon in coarse grain sizes. The microstructure in the SPD-processed lattice-softened Fe-Ni-Al-C specimens is still very complicated [1626], so further investigation on the effect of such microstructure on the Lüders deformation is required in the future.

## 7.6. High-entropy alloys

The newly developed generation of single-phase, multi-element (five
<figure><img src="image_156.png" /><figcaption>**Fig. 74.** (a, b) Mechanical properties and Lüders deformation behavior in the literature, where the number indicated at each plot corresponds to the reference number. (a) Total elongation versus ultimate tensile strength, (b) Lüders strain versus stress of Lüders deformation, $σ_L$. ●: Fe-Ni-Al-C [1602, 1603, 1605, 1608], ⚠: Fe-Mn [1604, 1606, 1607, 1611, 1612], △: Fe-Cr-Ni [1609, 1610], ×: Fe-Ni-Mn [1613]. (c) Stress-strain curves obtained by tensile tests for specimens rolled by 60% [1627]. (d) Local strain profiles along the tensile axis in the gauge section of the 5.0Al-60% cold-rolled specimen in (c).</figcaption></figure>

or more principal elements) solid solution alloys with concentrations between 5 and 35 at. percent of the individual elements, referred to as high-entropy alloys (HEAs) or multi-principal element alloys (MPEAs), has shifted the design concept away from the corners of phase diagrams toward the center, enabling compositions that go beyond the scope of traditional alloys [1638, 1639]. Such an approach offers unprecedented properties, challenges, and opportunities for a variety of structural and functional applications. HEAs exhibit certain remarkable mechanical properties. One of the most important aspects is the trade-off between strength and ductility, especially as the temperature decreases, resulting in excellent fracture toughness at cryogenic temperatures [1640, 1641].

Since SPD-processed materials with small grain sizes have been shown to have exceptional properties, including high yield strength and moderate ductility, and the ability to undergo superplastic forming [5], many scientists started to apply the SPD methods to HEAs. A current compilation of the publications covering SPD of HEAs is given in [1642]. The compilation shows that most of the work relates to FCC HEAs, mainly the equiatomic Cantor alloy CrMnFeCoNi [1643], a small part to BCC and dual phase FCC + BCC, but not to HCP HEAs. The inspection of publications also indicates that the main SPD process applied to HEAs is HPT. In the following sections, after reviewing the reported results about structural, microstructural and mechanical property changes in Cantor alloys after HPT processing, the hardening behavior, nanomechanical behavior, superplasticity and superficial properties of SPD-processed HEAs are reviewed.

### 7.6.1. *Microstructural and mechanical behavior of Cantor alloy*

Most SPD work has concentrated on the HPT processing of the Cantor

alloy [1642]. Therefore, this section aims to provide an overview of the HPT processing of this HEA, with emphasis on microstructure and texture evolution, phase transformation, strength and ductility as well as superplasticity and thermal stability [273]. During HPT processing of the Cantor alloy, a very fast refinement of the FCC microstructure takes place [290, 1644]. The grain size at room temperature and liquid nitrogen temperature (77 K) reaches a very low steady-state value of 24 nm after a shear strain of about 20 (Fig. 75a). Simultaneously, the dislocation density saturates at a high value of $3 \times 10^{16} \text{ m}^{-2}$ at room temperature (Fig. 75b), while it is surprisingly lower after cryogenic-SPD at 77 K ($10^{16} \text{ m}^{-2}$). The twin density at room temperature reaches a maximum value of 2% at a shear strain of 20 (Fig. 75c). During HPT at 77 K instead of mechanical twinning a deformation-induced phase transformation from FCC to HCP is observed (Fig. 75d) [290]. The onset pressure decreases with decreasing temperature [274], while there is a reverse transformation of the HCP phase produced by HPT at liquid nitrogen during long-term storage under ambient conditions [274]. The textures of HEAs processed by HPT are typical for shear, but they are quite weak [274, 290, 1645].

The room-temperature hardness of the Cantor alloy deformed by HPT at room temperature and in liquid nitrogen increases with shear strain and saturates at a certain strain level depending on pressure and HPT temperature (Fig. 76a) [274, 290]. Surprisingly, samples deformed in liquid nitrogen are softer (about 15%) than those deformed at room temperature. This also holds for compression down to 4 K (Fig. 76b) [291, 295]. The strength anomaly (softening) of the Cantor alloy is related to the martensitic phase transformation. Assuming that during pressure release and/or temperature increase as well as during
<figure><img src="image_157.png" /><figcaption>**Fig. 75.** (a) Crystallite size $\langle x \rangle_{area}$ (area-weighted mean crystallite size) [290], (b) dislocation density $\rho$ [290] and (c) twin density $\beta$ as a function of shear strain $\gamma$ for the Cantor alloy: after HPT processing at room temperature (open) and 77 K (full symbols) [290]. (d) Volume fraction of HCP martensitic phase versus applied pressure for the Cantor alloy after HPT processing at room temperature and in liquid nitrogen (LNT). The symbols represent different series of HPT samples, for which the phase composition was measured after the number of months indicated [274].</figcaption></figure>

room-temperature indentation, the HCP phase becomes unstable, then a reverse transformation is very likely. This process leads to a reduction of internal stresses and the formation of dislocation-free new grains of the FCC phase. Consequently, because of this process, the overall dislocation density of the FCC phase and the hardness of the polyphase aggregate are lowered. Moreover, the microstructure and microhardness become quite inhomogeneous [274]. The strength of the severely deformed Cantor alloy can be further increased by annealing (Fig. 76c) [503]. For this effect, precipitation seems to play a major role, along with other factors discussed in [689, 694, 702, 1646]. The softening above a certain temperature is due to the dissolution of precipitates and grain growth (precipitation leads to a loss of ductility) [503]. Moreover, the nanocrystalline Cantor alloy shows superplasticity above about half the melting temperature (Fig. 76d) [1647–1650]. This deformation mechanism is favored by reduced grain growth due to sluggish diffusion [1651] and/or nanosized precipitates often forming at elevated temperatures [1652, 1653].

In conclusion, similar to traditional metallic alloys, SPD processing of the Cantor alloy leads to UFG microstructures, high dislocation densities and weak textures. This is associated with a high strengthening effect and a loss of ductility. However, due to the multi-element solid solution, the Cantor alloy in the high-energy SPD state becomes quite unstable with respect to phase transformations, grain boundary segregation, and phase decomposition. Due to a defect-enhanced diffusivity [367], the kinetics of the decomposition is increased. This offers the possibility of producing ultra-hard materials by SPD processes or by post-annealing. Reduced grain growth, either caused by sluggish diffusion, segregation or nanoprecipitation, enables extremely high superplasticity at high strain rates [1650]. Although the effect of SPD on

changing properties has been demonstrated mainly for the Cantor alloy [273], a few examples of more complex HEAs indicate the high research potential of this new class of advanced materials in the field of SPD [1642, 1654–1656].

### 7.6.2. Hardening behavior of high-entropy alloys

The compositions of HEAs correspond to the unexplored middle parts of the phase diagrams; therefore, they exhibit never-before-seen behaviors such as high strength even at elevated temperatures, enhanced wear, oxidation, corrosion and radiation resistance [1638, 1643, 1657–1659]. As mentioned in the earlier chapter, the mechanical properties of these alloys can be further improved by applying SPD techniques. During SPD processing, both grain refinement and an increase in the density of lattice defects (e.g. dislocations and planar faults) contribute to the hardening of the as-processed materials [1, 2, 305, 855, 1659, 1660]. The maximum dislocation density and twin fault probability in HEAs achieved by SPD are about $300 \times 10^{14} \text{ m}^{-2}$ and 3%, respectively [1661–1665]. The minimum grain sizes for single-phase and multiphase HEAs were about 30 and 10 nm, respectively [1654, 1665]. These values are much smaller than the saturation grain sizes obtained in SPD-processed conventional metals and alloys as shown in Fig. 77, leading to high hardness of HEAs. The saturation state of the microstructure in HEAs is usually achieved at the equivalent strain of about 20 [1657, 1666].

The very high defect density and small grain size in SPD-processed HEAs cause superior hardness compared to conventional metals and alloys as shown in Fig. 78a [1654, 1661–1665]. On the other hand, the lower hardness bound of HEAs can also be achieved by conventional 316 L stainless steel processed by HPT at room temperature [19].
<figure><img src="image_158.png" /><figcaption>Fig. 76. (a) Room-temperature Vickers hardness as a function of shear strain γ after HPT processing at room temperature (open) and 77 K (full symbols) [290]. (b) Temperature dependence of the yield stress σ₀.₂ in nanograined Cantor alloy processed by HPT at room temperature and 77 K in comparison with coarse-grained sample [273]. (c) Vickers hardness of isochronally annealed (1 h) Cantor alloy samples after HPT processing at room temperature [503]. (d) Engineering stress-elongation curves at different temperatures for samples processed by HPT at room temperature [1647].</figcaption></figure>

Multiphase HEAs exhibit much higher hardness values than their single-phase counterparts [503, 1356, 1656, 1667–1677]. SPD-processed AlCrFeCoNiNb showed the highest hardness with a value of about 10 GPa due to the nanosized secondary phase precipitates [1654]. The hardening caused by SPD in the saturation state for different HEAs is shown in Fig. 78a where the hardness measured after SPD processing is plotted as a function of the values obtained before SPD. Fig. 78b reveals that SPD causes a higher relative hardening in HEAs

which had lower initial hardness values before SPD. Namely, below the initial hardness of about 3 GPa, the ratio of the hardness values after and before SPD varies between 1.6 and 4.5 while for HEAs with an initial hardness higher than 3 GPa, the hardening ratio is consistent with the value only of about 1.4. Fig. 78c shows the SPD-induced hardness increase versus the initial hardness in a double-logarithmic scale plot for different MPEAs having initial hardness values smaller than 3 GPa. The points in this plot follow an approximate linear trend which suggests the following relationship.

$$ \frac{H_{\text{SPD}}}{H_{\text{init}}} = A H_{\text{init}}^{\delta} \qquad (32) $$

where $H_{\text{SPD}}$ and $H_{\text{init}}$ are the hardness values given in MPa unit after and before SPD processing, respectively, and $A = 6870$ and $\delta = -1.05$ [1667].

As mentioned earlier, the high hardness of HEAs caused by SPD can be further enhanced by post-deformation annealing. This phenomenon was also observed for conventional SPD-processed alloys and called anneal hardening, as discussed earlier [689]. This effect was observed after annealing at moderate homologous temperatures between 0.3 and 0.4. Fig. 78d shows the hardness achieved by annealing after SPD versus the hardness observed immediately after SPD for different HEAs. The hardening effect of post-SPD annealing varies between 6% and 75%. Anneal hardening can be attributed to the reduction of mobile dislocation density, the relaxation of grain boundary structure and the formation of precipitates [1670, 1672, 1673, 1675–1677]. The highest hardness of about 11 GPa was achieved in AlNbTiV alloy processed by HPT at room temperature and then heat treated at 973 K [1677]. In this

<figure><img src="image_159.png" /><figcaption>Fig. 77. The grain size versus the dislocation density obtained for FCC and BCC Al-, Cu-, Ni- and Fe-based conventional alloys and multi-principal element alloys (MPEAs) processed by HPT at room temperature [1667].</figcaption></figure>
<figure><img src="image_160.png" /><figcaption>Fig. 78. (a) The maximum hardness measured after SPD processing versus the values obtained before SPD for different conventional materials and multi-principal element alloys (MPEAs). (b) The ratio of the hardness values measured after and before SPD as a function of the hardness determined before SPD for different conventional materials and MPEAs. (c) A double logarithmic plot of the data shown for MPEAs in (b). (d) The maximum hardness obtained due to anneal hardening of SPD-processed MPEAs versus the hardness values measured immediately after SPD [1667].</figcaption></figure>

alloy, anneal hardening was caused by precipitation. Due to superior hardness and strength, SPD-processed HEAs are expected to be used in different applications, e.g. as structural materials or surgical implants [1356, 1678].

### 7.6.3. Nanomechanical behavior of high-entropy alloys

Recently, the nanoindentation technique has been actively applied to characterize the mechanical behavior of HPT-processed medium-entropy alloys (MEAs) and HEAs [1679]. Moreover, employing different strain rates during nanoindentation has allowed the investigation of rate-dependent plasticity in HPT-processed alloys, providing valuable insights into the relationship between microstructural refinement and the enhancement of mechanical properties in these materials [1680]. This section describes the principles of nanoindentation for analyzing the rate-dependent plasticity in HEAs. The study of HEAs by nanoindentation was initially motivated by the relatively limited understanding of the grain refinement effect on rate-dependent plasticity in this new family of alloys compared to conventional materials.

The strain rate sensitivity (m) and activation volume (V*), which are two important indicators for rate-dependent plastic deformation, provide a clue to a better understanding of thermally-activated deformation mechanisms. According to the power-law response between nanohardness H_N, and indentation strain rate (˙e = h⁻¹(dh/dt), where h is displacement and t is time [1681]), the m and V* values can be determined at a given strain ε, and temperature T, by [1682]:

$$m = \left( \frac{\partial \ln H_N}{\partial \ln \dot{e}} \right)_{\varepsilon,T} = \left( \frac{\ln H_{N,2} - \ln H_{N,1}}{\ln \dot{e}_2 - \ln \dot{e}_1} \right)_{\varepsilon,T} \quad (33)$$

$$V^* = \sqrt{3kTC} \left( \frac{\partial \ln \dot{e}}{\partial H_N} \right)_{\varepsilon,T} = \sqrt{3kTC} \left( \frac{\ln \dot{e}_2 - \ln \dot{e}_1}{H_{N,2} - H_{N,1}} \right)_{\varepsilon,T} \quad (34)$$

where k is Boltzmann's constant and C is the constraint factor. Note that C is often taken as ~3, but it may vary with the material and loading conditions [1683]. In order to obtain m and V* values using nanoindentation, constant strain rate (CSR) tests and strain rate jump (SRJ) tests are typically employed. The examples of employing CSR and SRJ experiments on a CoCrFeMnNi HEA subjected to the HPT process are shown in Fig. 79. As shown in Figs. 79a and 79b, the HPT process produces a nanocrystalline structure with an average grain size of ~38 nm after 2 turns. In the case of CSR tests, multiple nanoindentation tests with different strain rates are performed and m and V* values are determined from the slope of a ln(H_N) versus ln(˙e) plot and the slope of the ln(˙e) versus H_N plot (Figs. 79c and 79d). On the other hand, for the SRJ test, ˙e is abruptly changed several times during a single nanoindentation (Figs. 79e and 79f). In order to account for the indentation size effect, H_N immediately before and after each transient change at specific depths is used for estimating m and V* [505, 1684]. These values are denoted as H_{N,1} and H_{N,2} in Equations 33 and 34, respectively.

To better understand the effect of grain refinement on the deformation mechanisms of M/HEAs, the m and V* values from the literature [1671, 1674, 1685–1704] are plotted as a function of grain size, d, in Figs. 80a and 80b. Note that only the results for FCC HEAs and MEAs are included in Fig. 80, since the variation in m and V* values with d in other crystal structures has not been extensively examined for these alloys. In the coarse-grain regime, the m values of HEAs and MEAs are
<figure><img src="image_161.png" /><figcaption>Fig. 79. Microstructures of (a) as-cast and (b) HPT-processed CoCrFeMnNi HEAs [1671] along with corresponding examples of (c,d) constant strain rate (CSR) nanoindentation [1671] and (e,f) strain rate jump (SRJ) nanoindentation [505].</figcaption></figure>

substantially higher than those for pure nickel, whereas HEAs, MEAs and pure nickel show comparable *m* values in the nanocrystalline regime. Compared to pure nickel, FCC HEAs and MEAs exhibit significantly lower *V*\* values (~10–100 *b*\*\*), indicating that the short-range of barriers (likely due to high lattice friction stress [1689, 1705] and chemical short-range ordering [1706, 1707]) are predominant. On the other hand, the *V*\* values for nanocrystalline FCC HEAs and nickel are around ~10 *b*\*\*, suggesting that the deformation mechanism in the nanocrystalline FCC HEAs is grain-boundary-mediated dislocation activity (e. g. dislocation nucleation and/or dislocation depinning at grain boundaries) as in the conventional nanocrystalline FCC metals [1696, 1697].

Fig. 80c shows the change in *m* of nanocrystalline CoCrFeMnNi HEA with testing temperature, suggesting that the *m* value increases abruptly from *T* = 473 K (~0.3 *T*<sub>m</sub>) [505, 1690]. To explore this mechanism,

transition, apparent activation enthalpy, Δ<i>Δ</i>\*, values can be determined by [1708, 1709]:

$$ \Delta\mathcal{H} = \frac{H_N V^* \frac{\partial(\ln H_N)}{\partial(1/T)}}{\sqrt{3C}} \quad (35) $$

Using Equation 35 and the *T*-dependent *H*<sub>N</sub> data (the inset of Fig. 80d), the variation in Δ<i>Δ</i>\* with temperature can be calculated (the main plot of Fig. 80d). The estimated Δ<i>Δ</i>\* for *T* ≤ 473 K is ~0.5 eV which indicates that the deformation is primarily governed by grain-boundary-mediated dislocation activities [1696, 1706, 1711]. On the other hand, the estimated Δ<i>Δ</i>\* for *T* > 473 K approaches that for grain boundary diffusion of nickel in CoCrFeMnNi HEA (~2.3 eV [1712]), suggesting that grain boundary diffusion becomes pronounced at higher temperatures.
<figure><img src="image_162.png" /><figcaption>**Fig. 80.** Variations in (a) strain rate sensitivity and (b) activation volume of various FCC metals and alloys as a function of grain size [1671, 1674, 1685–1704]. (c,d) Results of high-temperature nanoindentation experiments for HPT-processed nanocrystalline CoCrFeMnNi HEA [505, 1690].</figcaption></figure>

In summary, although appropriate for the estimation of strain rate sensitivity, activation volume, and apparent activation energy, nanoindentation provides crucial insights into the deformation mechanisms of SPD-processed HEAs. The results reported so far show that: (i) the deformation mechanism in FCC HEAs transforms from lattice atoms-dislocation interaction to gran-boundary-mediated plasticity as the grain size is reduced to the nanocrystalline regime, and (ii) grain boundary diffusion can become pronounced in nanocrystalline HEAs at homologous temperatures higher than ~0.3 $T_m$.

### 7.6.4. Superplasticity of high-entropy alloys

HEAs [1638, 1643, 1713–1720] and SPD processing [54, 159, 177, 193, 200, 224] are two extremes in materials science from the material and processing points of view. This resulted in the fast growth of research works on SPD processing of HEAs to achieve particular microstructures with advanced properties [5, 181, 1669, 1721, 1722]. Since superplasticity has a close connection with the SPD field [346, 770, 778, 783–786, 791], there have been attempts to examine the superplastic behavior of SPD-processed HEAs. The first evidence of superplasticity in HEAs was reported more than ten years ago in a fine multiphase equiaxed structure of an AlCoCrCuFeNi alloy with an average grain size of ~2 µm processed through MDF at 1223 K [1723, 1724]. In these experiments, a maximum elongation of 1240% was achieved when testing at 1273 K at a strain rate of 1.0 × 10⁻² s⁻¹. More recently, there have been several reports of high superplastic elongations in HEAs and a detailed tabulation of all results is presented elsewhere [1725]. One of the most striking results available to date includes an elongation of 2000% in a CoCrFeNiMnAl₀.₅ alloy [1650]. Fig. 81 shows results obtained on a Mo₇.₅Fe₅₅Co₁₈Cr₁₂.₅Ni₇ HEA where tensile tests were conducted at 973 and 1273 K [1726]. As shown in Fig. 81a, this alloy revealed the presence of a rather complex multiphase

microstructure including FCC, BCC, martensite and an Mo-rich µ phase before tensile testing. The results in Fig. 81b demonstrate the occurrence of superplastic elongations of 500% when testing at 1273 K but with an absence of superplasticity at 973 K. The latter result was attributed to the formation of more Mo-rich µ precipitates at the higher testing temperature. Close inspection of the superplastic samples in Fig. 81b shows that they represent true superplasticity because there is no evidence of any additional necking within the gauge lengths [1727].

An important requirement in superplasticity is to test the applicability of the grain boundary sliding model (i.e. general creep equation Equation 7) which was first developed for conventional superplastic alloys without processing by SPD techniques. Several analyses demonstrate that there is very good agreement between published data for various conventional Al- and Mg-based alloys and the predictions of the general creep equation for superplasticity [1728–1730]. The results of this analysis are shown in Fig. 82 where the temperature and grain size compensated strain rate is plotted against the normalized stress. The data for Fig. 82 were taken from several published reports [1481, 1647, 1650, 1652, 1653, 1724–1726, 1731–1734] and full details of the analysis were given elsewhere [1725]. The solid line labeled $\dot{\epsilon}_{sp}$ represents the prediction of the model and it is apparent that there is generally excellent agreement between the experimental datum points and the rates predicted by the superplastic model based on the occurrence of grain boundary sliding [791]. This analysis confirms that superplasticity in HEAs is similar to the superplasticity occurring in other materials both with and without SPD processing.

### 7.6.5. Superfunctional high-entropy alloys and ceramics

High-entropy materials include two major groups of HEAs and high-entropy ceramics. Similar to HEAs [1735, 1736], high-entropy ceramics are defined as materials with five or more cations or with an entropy of
<figure><img src="image_163.png" /><figcaption>**Fig. 81.** (a) XRD patterns of Mo<sub>7.5</sub>Fe<sub>55</sub>Co<sub>18</sub>Cr<sub>12.5</sub>Ni<sub>7</sub> HEA before (annealed) and after HPT processing [1650]. (b) Elongation versus strain rate plot for HPT-processed Mo<sub>7.5</sub>Fe<sub>55</sub>Co<sub>18</sub>Cr<sub>12.5</sub>Ni<sub>7</sub> before and after tensile testing at elevated temperatures [1727].</figcaption></figure>

mixing more than 1.5R (R: gas constant) (Fig. 83a) [986]. The high entropy of these materials results in low Gibbs free energy and their consequent high stability [983,1732,17333]. High-entropy materials have shown high potential for a varied range of not only mechanical properties but also functional properties due to their superior features including sluggish diffusion, lattice strain and the cocktail effect [1737–1744]. SPD methods, especially HPT, have been employed for processing and synthesizing high-entropy materials mainly for mechanical properties [391, 1236, 1745–1749] and partly for achieving functional properties [1145]. Ultrahigh hardness, hydrogen embrittlement resistance, hydrogen storage capability, biocompatibility, photovoltaic and photocatalysis are the reported functional properties and applications of SPD-processed/synthesized high-entropy materials, which are briefly reviewed below [1642].

* In connection with classic works on SPD processing of conventional materials [4, 38], there have been attempts to improve the mechanical properties of HEAs by SPD methods to levels comparable to those of ceramics [1642]. A combination of the concept of SPD and HEAs resulted in the introduction of the two HEAs processed by HPT with ultrahigh hardness levels up to 10 GPa [1654, 1655]. Among all reported HEAs, carbon-doped AlTiFeCoNi as a dual-phase HEA, which was designed by CALPHAD (calculation of phase diagram)

<figure><img src="image_164.png" /><figcaption>**Fig. 82.** Temperature and grain size compensated strain rate for SPD-processed HEAs with superplasticity behavior, showing excellent agreement with the theoretical prediction for conventional superplasticity [1650]. $\epsilon_{sp}$: strain rate, k: Boltzmann's constant, T: absolute temperature, $D_{gb}$: coefficient for grain boundary diffusion, d: grain size, b: Burgers vector, $\sigma$: applied stress G: shear modulus). Experimental data were taken from [1481, 1647, 1650, 1652, 1653, 1724–1726, 1731–1734].</figcaption></figure>

method and processed by HPT, showed an ultra-high hardness of 950 Hv [1655]. One of the highest hardness values (1030 Hv) was achieved in HPT-processed AlCrFeCoNiNb HEA with six principal elements and two cubic and hexagonal phases (Fig. 83b) [1654]. These high levels of hardness values of HEAs are attributed to the formation of nanograins, dislocation and nanoprecipitates [1642, 1654].

* As mentioned earlier, one critical issue in developing a hydrogen-based energy system is achieving materials with high hydrogen embrittlement resistance [1750]. Enhanced hydrogen embrittlement resistance was observed in some FCC HEAs due to their high plasticity in the presence of hydrogen [1240, 1751], but coarse-grained HEAs suffer from low yield strength. It was reported that HPT was effective in enhancing the yield strength of CrMnFeCoNi HEA as a hydrogen-compatible material, as shown in Fig. 83c [1235, 1752]. Dark-field images and high-resolution TEM images of CrMnFeCoNi after HPT processing show the presence of defects with low mobility such as twins (Fig. 83d), Lomer-Cottrell locks (Fig. 83e) and D-Frank partial dislocations (Fig. 83d) which should be responsible for suppressing hydrogen-enhanced localized plasticity (HELP) mechanism [1752].

* Solid-state hydrogen storage at room temperature is another critical issue in developing a hydrogen-based energy system [1076, 1753]. For the solid hydrogen storage concept, since the SPD methods such as ECAP and HPT could improve the properties of conventional metal hydride in terms of kinetic and thermodynamic [359, 1754], there were some attempts to synthesize new HEAs such as MgTiVCrFe by HPT for hydrogen storage [314]. Although MgTiVCrFe showed low hydrogen storage capacity, this study was a starting point to develop various HEAs for fast and reversible hydrogen storage at room temperature by employing theoretical studies (Fig. 84a): TiZrCrMnFeNi (Fig. 84b) [1755], TiZrNbFeNi (1756), TiZrNbCrFe [1757] and Ti<sub>x</sub>Zr<sub>2-x</sub>CrMnFeNi [1758] (x = 0.4–1.6). These materials
<figure><img src="image_165.png" /><figcaption>Fig. 83. (a) Entropy of mixing versus number of elements [986]. (b) Hardness of AlCrFeCoNiNb synthesized by HPT for N = 1/16, 1/4, 1 and 4 turns [1654]. (c) Engineering stress versus engineering strain before and after hydrogen exposure for CrMnFeCoNi before and after HPT processing with varied turns [1752], (d) Dark-field image of twins with relevant selected area electron diffraction patterns for CrMnFeCoNi after HPT processing for 1/16 [1752]. (e,f) High-resolution lattice images of CrMnFeCoNi after HPT processing for 1/4 turn containing (e) Lomer-Cottrell locks and (f) D-Frank partial dislocations [1752].</figcaption></figure>

also showed a promising potential to be used as anode materials in nickel-metal hydride batteries [1759].

* Biocompatibility is another application of HEAs, which is getting high attention [1760, 1761]. Application of the HPT method to a material designed by CALPHAD calculations (TiAlFeNiCo HEA) was reported to be effective in enhancing the strength, decreasing the elastic modulus (by phase transformation) and improving the biocompatibility of this material [1356]. The material showed a larger microhardness and better biocompatibility (cell proliferation-viability-cytotoxicity activity) compared to coarse-grained and UFG pure titanium and the Ti-6Al-7 Nb alloy representing two conventional biomaterials (Fig. 84c) [1356]. In another report, ultra-SPD with extremely high shear strain was employed to synthesize TiNbZrTaHf HEA with BCC structure as a biomaterial [317]. The material showed higher hardness compared to corresponding binary, ternary and medium-entropy alloys with an elastic modulus better than that of some typical biomaterials [317].

* Photovoltaic and photocatalysis are two functional properties of some high-entropy ceramics. The semiconductors having these two properties can adsorb the light to generate electricity (photovoltaic) and consequently advance a reaction (photocatalysis). TiZrHfNb-TaO₁₁ high-entropy oxide [323, 324, 982] and TiZrHfNbTaO₆N₃ high-entropy oxynitride (HEON) [326, 327, 982] are two high-entropy ceramics synthesized by HPT which could generate the electric current under light as shown in Fig. 84d. The potential of HPT to improve the photovoltaics was also reported in conventional semiconductors [963]. These two high-entropy ceramics also showed activity for photocatalytic hydrogen evolution and the

efficiency of TiZrHfNbTaO₆N₃ was higher than Ga₆ZnON₆ as shown in Fig. 84e [323, 326]. High-resolution TEM image of TiZrHfNb-TaO₆N₃, which contains FCC and monoclinic phases and has higher light absorbance than TiZrHfNbTaO₁₁ due to its low bandgap, is shown in Fig. 84f [326, 327]. As mentioned in the earlier chapter, these two materials later were employed for photocatalytic CO₂ conversion with high efficiency [324, 327]. TiZrHfNbTaO₁₁ showed a high CO production rate the same as P25 TiO₂ used as a benchmark photocatalyst [324]. The efficiency of TiZrHfNbTaO₆N₃ for CO₂ photoreduction was even higher than that of P25 TiO₂ and all reported photocatalysts in the literature [327]. TiZrNbTaWO₁₂ is another HEO synthesized by arc melting and processed by HPT which showed photocatalytic activity for oxygen evolution [325] due to the heterojunction effects [1762].

The combination of the concepts of high-entropy materials and SPD resulted in the discovery of various superfunctional high-entropy materials for various ranges of applications. Perhaps, the introduction of the first high-entropy photocatalysts is the biggest contribution of the SPD field in this regard. Since there are urgent needs for functional biomaterials and energy materials for carbon-neutral energy development, it is expected that the application of SPD in the field of high-entropy materials will be extended from mechanical properties to functional properties in the near future.

### 7.7. Amorphous and glass materials

Amorphous alloys are one of the most attractive topics of modern
<figure><img src="image_166.png" /><figcaption>**Fig. 84.** (a) Lattice structure of Ti<sub>0.4</sub>Zr<sub>1.6</sub>CrMnFeNiH<sub>6</sub> high-entropy hydride constructed by first-principles calculations [1758]. (b) Hydrogen storage pressure-composition isotherms at room temperature for TiZrCrMnFeNi showing reversible hydrogen storage [1755]. (c) MTT cell viability assay (biocompatibility) and microhardness of TiAlFeCoNi alloy before and after HPT processing compared to titanium and Ti-6Al-7 Nb biomaterials [1356]. (d) Photocurrent generation of TiZrHfNbTaO<sub>11</sub> versus time [324]. (e) Photocatalytic hydrogen production rate of TiZrHfNbTaO<sub>6</sub>N<sub>3</sub> versus time compared to convention oxynitride Ga<sub>6</sub>ZnON<sub>6</sub> [326]. (f) High-resolution TEM image of TiZrHfNbTaO<sub>6</sub>N<sub>3</sub> containing nanograins of two phases [326].</figcaption></figure>

materials science [1763–1765]. They are usually obtained as thin ribbons by the rapid quenching from the melt at typical cooling rates of about 10⁶ K/s [1763] or in the form of bulk metallic glass (BMG) in the shape of cylinders with diameters of up to several centimeters at lower cooling rates of about 10² K/s [1764]. Deformation of amorphous alloys occurs through the formation and propagation of shear bands with a small thickness (~10 nm). Such behavior leads to extremely low tensile ductility; however, amorphous alloys exhibit some ductility in bending and compression tests. It is possible to improve the ductility of amorphous alloys by increasing the so-called free volume and by creating heterogeneities in the amorphous phase [1765]. A promising way to introduce high strain and therefore, to transform the structure of the glassy metals is the use of SPD through the HPT process [1766]. There are many research groups actively involved in studying the effect of HPT treatment on amorphous alloys and a few reviews were published as well [1087, 1767]. The results obtained can be summarized as follows.

HPT allows the consolidation of amorphous rapid-quenched ribbons and powders into monolithic samples [1768–1770]. The structural

changes in amorphous alloys during HPT are determined by their chemical composition. For instance, HPT treatment of Al-rich [1770], Fe-rich [529] and Nd-Fe-B alloys [1771] leads to nanocrystallization in the amorphous phase. In the Ti₅₀Ni₂₅Cu₂₅ rapid-quenched alloy TiNi nanocrystals of the B2 phase precipitate in the amorphous phase [1771, 1772]. An attempt to analyze these processes has been reported in [1767, 1773]. In amorphous alloys without a significant excess of the basic element in the composition, for example, in Zr-based BMG, nanocrystallization by HPT is not observed in most cases, but HPT leads to an increase in the free volume and the enthalpy of the amorphous phase (Fig. 85a) [1772, 1774, 1775]. In some cases, a cluster-type structure was formed, which is associated with the chemical separation in the amorphous phase as a result of HPT processing [1776, 1777]. Chemical separation along the shear band in HPT-treated Zr-based amorphous alloy was unveiled by advanced microscopy methods [1778].

Properties of amorphous alloys change significantly by HPT processing, in particular, the microhardness decreases (Fig. 85b) [1774,