Dawn of nitride ferroelectric semiconductors: from materials to devices

III-nitride semiconductors are promising optoelectronic and electronic materials and have been extensively investigated in the past decades. New functionalities, such as ferroelectricity, ferromagnetism, and superconductivity, have been implanted into III-nitrides to expand their capability in next-generation semiconductor and quantum technologies. The recent experimental demonstration of ferroelectricity in nitride materials, including ScAl(Ga)N, boron-substituted AlN, and hexagonal BN, has inspired tremendous research interest. Due to the large remnant polarization, high breakdown field, high Curie temperature, and significantly enhanced piezoelectric, linear and nonlinear optical properties, nitride ferroelectric semiconductors have enabled a wealth of applications in electronic, ferroelectronic, acoustoelectronic, optoelectronic, and quantum devices and systems. In this review, the development of nitride ferroelectric semiconductors from materials to devices is discussed. While expounding on the unique advantages and outstanding achievements of nitride ferroelectrics, the existing challenges and promising prospects have been also discussed.


Introduction
One of the recent significant developments of III-V compound semiconductors and ferroelectrics is the realization of III-nitride (III-N) ferroelectric semiconductors: It has been shown, both theoretically and experimentally, that the addition 1 These authors contributed equally.
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Original content from this work may be used under the terms of the Creative Commons Attribution 4.0 licence. Any further distribution of this work must maintain attribution to the author(s) and the title of the work, journal citation and DOI. of rare-earth metal elements into III-nitride lattice, such as scandium (Sc) in AlN, can flatten the Gibbs free-energy landscape of the III-nitride wurtzite structure, transforming conventional III-nitrides to be ferroelectrics [1][2][3][4][5]. Some of the major historical developments of nitride ferroelectric materials are illustrated in figure 1. Previously, ferroelectricity can only be observed in complex crystal structures like oxide perovskites, whose performance suffers from the presence of extensive oxygen vacancies [6,7]. Compared to oxide ferroelectrics, nitride ferroelectrics are completely free of oxygen vacancies and are also with relatively less nitrogen vacancies due to the stronger metal-nitride bond and the higher nitrogen vacancy formation energy [8], thereby promising high stability and endurance [9,10]. Moreover, nitride ferroelectrics exhibit unprecedentedly high paraelectric transition temperature (Curie temperature), linear displacement, large remnant polarization, tunable coercive field, and compatibility with mainstream semiconductor manufacturing process, such as Si, GaN, as well as complementary metal-oxide semiconductor (CMOS) and integrated circuits technologies [5,[11][12][13][14][15].
Compared to conventional III-nitride semiconductors, studies have shown that nitride ferroelectric semiconductors, such as ScAlN, can exhibit significantly enhanced electrical, piezoelectric, as well as linear and nonlinear optical properties. For example, the piezoelectric coefficient d 33 and permittivity of Sc 0.4 Al 0.6 N are nearly five and two times larger than that of AlN, respectively [16][17][18][19]. Recent studies further suggest that ScAlN possesses unusually large optical χ (2) nonlinearity, which was measured to be at least one order of magnitude higher than AlN and twice the value of the extensively studied LiNbO 3 [20,21]. These unique characteristics, together with its ultra-wide bandgap, ferroelectric functionality, and seamless integration with III-nitride technology, have made ScAlN one of the most promising materials for a wide range of future applications, including high power and high frequency electronics [22][23][24], acoustic resonators and filters [25][26][27], micro-electromechanical systems [28][29][30], memory electronics [13,31,32], neuromorphic computing [33,34], quantum photonic circuits [35], and quantum transduction from microwave to infrared, visible and ultraviolet (UV) [36][37][38][39], to name just a few (figure 2).
Ferroelectric ScAlN was first demonstrated by sputter deposition [5], which, however, offers very limited material quality, leading to non-optimal device performance. Recently, it has been shown that single crystalline ferroelectric ScAlN can be realized utilizing standard epitaxial techniques, e.g. molecular beam epitaxy (MBE) [10], offering for the first time the ability to truly unlock the potential of this material system. The fully epitaxial growth of Sc-III-N provides a viable path to integrate the piezoelectric, ferroelectric, optical, and quantum functionality with well-established semiconductor platforms and further promises a thickness scaling into the nanometer or even monolayer regime.
Despite rapid progress made in nitride ferroelectric semiconductor materials and devices, to date, there has been no systematic review articles on this emerging topic to our knowledge. In this article, we therefore provide a detailed overview of the major advances of nitride ferroelectric semiconductor materials and a broad range of emerging device applications. In section 2, the synthesis/growth of ScAlN and ScGaN is described. State-of-the-art material properties utilizing various growth techniques, including sputter deposition, MBE, and metalorganic chemical vapor deposition (MOCVD) are also discussed and compared. Special attention is paid to the control of lattice-polarity and impurity incorporation. In section 3, the fundamental material properties, including energy bandgap, dielectric properties, piezoelectricity, ferroelectricity, and linear and nonlinear optical properties are discussed. Their emerging device applications, including high electron mobility transistors (HEMTs), memory electronics, photodetectors, acoustic resonators and filters, and quantum photonics, are presented in section 4. Finally, summary and prospects are described in section 5.  [5, 14, 16, 20, 43, 60-62, 64, 123, 145]. b [10,15,19,24,32,45,47,48,69,91,106,125,217,225]. c [51][52][53]. d (0002) plane XRC FWHM for Sc 0.2 Al 0.8 N. e The lowest impurity concentration was achieved with a specific growth strategy [69].

Material growth
ScN has a stable ground-state rock-salt phase (space group Fm3m, point group O h , nonpolar) [40], which is not compatible with the stable wurtzite structure (space group P6 3 mc, point group C 6v , polar) of conventional III-nitrides [41]. However, Takeuchi reported in 2002 that an additional local minimum of total-energy existed in wurtzite ScN based on the first-principle calculations, making it possible for the synthesis of wurtzite phase Sc-III-N alloys [42]. Using localdensity approximation calculations, Farrer and Bellaiche predicted that ScN has a metastable hexagonal structure instead of the wurtzite structure in the same year [1]. In 2013, Zhang et al studied the evolution of crystal structure of Sc-III-N using density functional theory (DFT), and reported that ScAlN and ScGaN alloys have a stable wurtzite phase with Sc contents up to about 0.56 and 0.66, respectively, above which rocksalt phases are more favorable [4]. Obviously, growth of pure wurtzite phase high quality Sc-III-N alloys with high Sc contents is challenging, but wurtzite phase Sc-III-N is essential for realizing the unprecedented ferroelectricity in III-nitride semiconductors [2,3]. A variety of growth techniques have been utilized over the past decades for Sc-III-N deposition, including reactive magnetron sputtering deposition [16,17,43,44], MBE [45][46][47][48], MOCVD [49][50][51][52][53], and hydride vapor phase epitaxy [54][55][56]. In this part, we will focus mostly on sputter, MBE, and MOCVD for high-quality Sc-III-N thin films growth. Table 1 briefly summarizes some of the material properties of ScAlN grown by these techniques.

Sputtering deposition of Sc-III-N alloys
To date, the majority studies on Sc-III-N alloys have employed sputtering deposition techniques for the material preparation. Various sputtering deposition conditions, including deposition temperature [17], power [57], pressure [58], and atmosphere (N 2 /Ar flow ratio) [59], and substrates, such as Pt [5], Mo [60], Al [14], sapphire [43], Si [16], MgO [43], etc, have been widely studied to improve the phase purity, crystallinity, crystalline quality, grain size, built-in stress, surface morphology, resistivity, and piezoelectric response of ScAlN films. In 2009, Akiyama et al reported the ScAlN solid solution films deposited at 580 • C on Si(100) substrates using a dual radio frequency (RF) magnetron reactive co-sputtering, in which solely wurtzite phase, coexistence of wurtzite and cubic phases, and cubic phase were observed for the Sc contents less than 0.41, between 0.42 and 0.45, and higher than 0.46, respectively [16]. However, they observed a dramatic degradation of wurtzite phase accompanied with the degeneration of crystal quality as well as the piezoelectric coefficient for the Sc 0.33 Al 0.67 N film. In their following studies, they found that better crystalline quality and higher piezoelectric response can be achieved by lowering the deposition temperature to 400 • C, which was attributed to highly ordered nanograin growth [17].  [44]. Despite great efforts have been made to improve the crystallinity and crystal quality, sputtered ScAlN films generally have a highly c-axis oriented texture structure and an x-ray diffraction (XRD) rocking curve (XRC) full-width-at-half-maximum (FWHM) larger than 3600 arcsec [61][62][63]. Most recently, Schönweger et al sputtered ScAlN films on commercial Si-doped GaN/sapphire templates, however, the (0002) plane XRC FWHM for a 200 nmthick Sc 0.28 Al 0.72 N film was still as high as 3600 arcsec [64]. Meanwhile, by detailed annular bright field scanning transmission electron microscope analysis (figure 3(c)), they found that the lattice-polarity was changed from the initial metal (M)-polar to a multi-domain state, i.e. M-polar and nitrogen (N)-polar coexisted, after the ScAlN film thickness is beyond 30 nm [64]. This polarity transition was attributed to the competition between a M-polar GaN template and a deposition process that favors N-polar growth [64]. Therefore, precisely control the as-grown lattice-polarity of sputtered Sc-III-N films remains challenging. Sputtered ScGaN films have been also reported by several groups since 2001, and the growth behaviors and microstructures are similar to sputtered ScAlN [65][66][67][68].

MBE
Epitaxial growth of crystalline ScN has been studied in early 1990s [76,77], while alloying Sc with nitrides under the wurtzite framework was achieved in early 2000s. In 2004, Constantin et al reported the growth of ScGaN using a RF MBE over the Sc fraction range of 0-1 under N-rich conditions [78]. The growth of ScGaN was separated into three regimes: (a) wurtzite like but having local lattice distortions for small Sc contents (less than 0.17), (b) a transitional region with mixed phases for intermediate Sc contents (0.17-0.54), and (c) a cubic, rock-salt like for large Sc contents (higher than 0.54). It was found that the addition of Sc atoms caused local lattice distortions in wurtzite phase ScGaN [79]. In 2009, Moram et al reported the growth of wurtzite ScGaN with low Sc content (up to 0.08) using a NH 3 -MBE under N-rich conditions [80]. Later, it was found that the growth rates of ScGaN decreased significantly compared to GaN grown under similar conditions, which was attributed to the presence of a Sc surfactant layer on the surface [81]. In 2015, Tsui et al reported the growth of ScGaN under metal-rich conditions using a RF MBE, and realized phase-pure wurtzite ScGaN with a Sc content up to 0.26, which is significantly higher than that previously reported for N-rich conditions [82]. Additionally, it was found that the growth rate increased with increasing Sc flux, which was interpreted to the catalytic decomposition of N 2 induced by the presence of Sc atoms [82]. In 2021, Wang et al reported the growth of wurtzite ScGaN with a Sc content up to 0.41 by using a RF MBE system, and further demonstrated, for the first time, ferroelectricity in ScGaN [83].
However, the growth of phase-pure wurtzite ScAlN using MBE has been only achieved recently, due to the coexistence of multiple phases in the Sc-Al-N phase diagram, including ScAl intermetallic, perovskite phase Sc 3 AlN 1−x , and wurtzite and cubic phase ScAlN [84]. In 2017, Hardy et al first reported the growth of ScAlN on GaN and SiC substrates using MBE [45]. It is well known that for the growth of conventional IIInitrides using MBE, slightly metal-rich conditions are beneficial to improve the surface morphology and crystal quality [85][86][87][88][89]. However, when ScAlN is grown under metal-rich conditions, excess Sc segregates to the surface due to the preferred incorporation of Al, resulting in mixed phases, degenerated crystal quality, and rough surface, whereas excellent film quality and pure phase were demonstrated under N-rich growth conditions [45,46,90]. Figure 4(a) shows the (0002) plane XRD 2θ−ω scans for 200 nm-thick ScAlN films grown on SiC substrates at 700 • C with varying III/V ratios [90]. For a III/V = 0.3-0.9, a single diffraction peak for wurtzite ScAlN was observed. In contrast, for III/V = 1, three separate peaks at 36.1 • , 36.8 • , and 37.9 • were observed, indicating the presence of mixed phases [90]. Most recently, Engel et al proposed a metal-modulated epitaxy (MME) approach to grow ScAlN under metal-rich conditions using MBE, wherein phase-pure ScAlN films with high crystal quality and smooth surface were also achieved [91].
Several approaches have been reported to improve the phase purity of MBE-grown ScAlN [46,90]. In 2019, Frei et al reported the growth of wurtzite phase ScAlN by pulsed supply of Sc and Al under N-rich conditions [46]. In 2020, Hardy et al reported that a reduction in growth temperature from 700 to 390 • C was required for Sc 0.32 Al 0.68 N to maintain the high quality wurtzite structure [90]. This result suggests that ScAlN has a wide growth temperature window. Figure 4(b) displays the (0002) plane XRD 2θ−ω scans for Sc 0.2 Al 0.8 N/GaN samples grown over a temperature range of 400 • C-900 • C under N-rich conditions [47]. The characteristic diffraction peak for wurtzite crystal structure was shown in all samples, and gradually shifted toward the higher angle side with increasing growth temperature, due to the residual tensile strain induced by the thermal expansion coefficient mismatch [92,93]. With these efforts, high quality wurtzite phase ScAlN with a Sc content up to 0.4 has been achieved utilizing MBE [46][47][48]. Figures 4(c) and (d) show the reflection high energy electron diffraction (RHEED) patterns [47] and atomic force microscope (AFM) images [94] of MBE-grown ScAlN epilayers with various Sc contents. Generally, the N-rich growth conditions result in a granular surface. The nanograins on the surface have a size of a few nanometers to tens of nanometers, which increases with increasing Sc content [47,94]. Obviously, the crystallinity and crystal quality of MBE-grown ScAlN, especially ScAlN with higher Sc contents, are dramatically improved compared to sputtered films.

MOCVD.
MOCVD has been commonly utilized to grow high quality III-nitride films, heterostructures, and device structures with a relatively higher growth temperature and larger growth rate, which has brought III-nitride materials from fundamental research to industrial production [95][96][97]. To date, however, the growth of Sc-III-N alloys by MOCVD has remained very challenging. In 2009, Koleske et al reported issues associated with MOCVD growth of ScGaN [49]. Due to the low volatility of Sc metalorganic precursor, only doping level concentration of Sc was achieved in GaN film. In 2013, Saidi et al also reported the growth of Sc-doped GaN, and they found that the coalescence process and near band edge emission of GaN was significantly influenced by the Sc incorporation [50]. Until 2019, Leone et al first reported the deposition of ScAlN with Sc content up to 0.3 by MOCVD [51]. For MOCVD growth, choosing a proper precursor for Sc is the first step. To let Sc enter the MOCVD reaction chamber and incorporate into the epilayer as a main element with alloy level, a Sc precursor with high vapor pressure is necessary. On the other hand, to avoid the incorporation of oxygen, the Sc precursor should be free of oxygen [51]. With these restrictions, they finally chose commercially available scandiocene tris (methylcyclopentadyenil)Sc (Cp 3 Sc) as Sc precursor [51]. Cp 3 Sc was used by Saidi et al before, however, only doping level incorporation of Sc was achieved, due to the low pressure of Sc precursor when running the Cp 3 Sc bubbler in the normal conditions (heated to 90 • C and held at atmospheric pressure) [50]. Leone et al found that Cp 3 Sc bubbler should be heated up to 150 • C to reach a flow that is comparable with the flow of trimethylaluminum used for AlN deposition [51]. Unfortunately, this temperature (150 • C) exceeds the temperature limitation of most of the sealing and electronic components equipped in a MOCVD system. Therefore, they upgraded their MOCVD system to satisfy the temperature requirement for the Cp 3 Sc bubbler, and achieved the growth of wurtzite-type ScAlN on sapphire and SiC substrates using a GaN buffer layer [51][52][53].

Lattice-polarity control.
In the wurtzite structure, III-nitride semiconductors, i.e. AlN, GaN, InN, and their alloys, possess a strong spontaneous polarization along caxis [41,98,99]. Therefore, two stabilized states with antiparallel polarization direction exist, denoted as M-polar and N-polar. As polarization engineering plays a crucial role in both epitaxial growth and device design for III-nitrides, especially for nitride ferroelectric semiconductors, there is an urgent need to precisely control the lattice-polarity of Sc-III-N  [24] with the permission of AIP Publishing). [24,48,[100][101][102][103]. However, the widely employed sputter technique for Sc-III-N deposition generally favors an N-polar as-grown lattice [5,60,64]. Early studies have shown that MBE has the unique ability to control the lattice-polarity of III-nitrides through buffer layer, insert layer, interfacial layer, as well as doping engineering [100,103,104]. Yet to date, most of the MBE-grown Sc-III-N films exhibited an M-polar lattice [45,47,48]. In 2021, Wang et al reported the growth of phase-pure wurtzite N-polar ScAlN on sapphire substrates by controlling its lattice-polarity through the underlying GaN buffer layer [24].  [24,47]. XRD reciprocal space mapping for the asymmetric (105) plane confirmed a coherent growth for the lattice-matched Sc 0.18 Al 0.82 N on GaN ( figure 5(b)). Similar to other III-nitrides, the latticepolarity of ScAlN can also be determined by wet chemical etching. After etching in TMAH solution, pyramidal nanostructures were observed on the N-polar ScAlN surface, confirming the N-polar lattice (figures 5(c) and (d)).

Impurity control.
Due to the large electronegativity difference between Sc and oxygen, Sc has a large oxygen affinity like Al, which therefore leads to a severe oxygen incorporation issue during ScAlN growth [94,105]. On the other hand, it is difficult to purify the rare-earth (RE) elements due to their special chemical properties. As a result, the commercially available Sc sources generally have a lower purity (3-5 N) than other IIIA metals (e.g. Al, Ga, and In, which have a purity of 6-7 N). In 2020, Casamento et al reported an oxygen concentration of 10 19 -10 21 cm −3 for ScAlN with a Sc content up to 0.26 grown on AlN template under N-rich conditions using a Sc source with a purity of 4N [105]. Most recently, they compared the effect of Sc source purity on the impurity incorporation during ScAlN growth and leakage during electrical measurements [106]. It was found that a higher purity source leads to a 2-3 orders of magnitude reduction in the oxygen, carbon, and fluorine impurity concentration in ScAl-N/GaN multilayers, resulting in a 5-7 orders of magnitude reduction in the electrical leakage in the ScAlN/GaN single heterostructure.
The photoluminescence (PL) of ScAlN is also extremely sensitive to the impurity incorporation. Figure 6(a) shows the PL spectra of Sc 0.05 Al 0.95 N grown on AlN template using MBE [94]. AlN has a dominant free exciton emission at 210 nm. However, no emission was observed in the predicted direct bandgap range (5.7-6.1 eV) for Sc 0.05 Al 0.95 N. Instead, a broad and bright defect-related emission at 352 nm was measured. Through detailed power-and temperaturedependent PL measurements and analysis, the broad peak was attributed to oxygen defect dominated emission, including charge carrier recombination from the localized excited state of (V cation -O N ) 2−/− to its ground state, and charge carrier recombination of isolated V 3−/2− cation to valance band. This oxygen defect dominated emission can be significantly suppressed by increasing growth temperature. The suppressed incorporation of oxygen and carbon impurity at elevated growth temperatures has been also reported in MOCVDgrown ScAlN films [52]. In addition, the oxygen defect dominated emission weakened gradually with the increase of Sc content, due to the increasing nonradiative recombination process induced by high impurity concentration (figure 6(b)) [94].
Most recently, Wang et al proposed a Ga-assisted growth strategy to improve the material quality and suppress the impurity incorporation of ScAlN by forming ScAlGaN quaternary alloy [69]. With this Ga-assisted growth method, the commonly used N-rich condition was well compensated by the excess Ga and the migration ability of impinging atoms was significantly enhanced as well. Normally, ScAlN films grown with N-rich conditions exhibited a granular surface, while smooth surface with spiral atomic steps were observed for ScAlGaN epilayers grown with Ga-rich conditions. The highly ordered atomic stacking sequence shown in the high-angle annular dark field STEM images of ScAlGaN confirms a single crystalline wurtzite structure (figures 7(a)-(c)). Furthermore, the Ga-assisted growth strategy also showed advantages in achieving a clean and sharp interface (figure 7(d)), which was evidenced by the clear periodicity in ScAlGaN/GaN superlattices, and the improved 2DEG mobility in ScAlGaN/GaN HEMT structures. The XRC FWHM for both (0002) and (1012) planes of ScAlGaN was drastically reduced compared to ScAlN with a similar Sc content (figure 7(e)). Moreover, the Ga-assisted growth method demonstrated a promising path to suppress the impurity incorporation in ScAlN. Figure 7(f) displays the secondary ion mass spectrometry (SIMS) profile recorded in a ScAlGaN/GaN multilayer stack. The average oxygen and carbon concentration in all ScAlGaN layers (with a relative Sc content over a range of 0.17-0.30) was below 3 × 10 17 and 5 × 10 16 cm −3 , which is 3-4 orders of magnitude lower than that of ScAlN with similar Sc contents grown on AlN template using the same Sc source (purity 3 N). These impurity concentration is also lower than the ScAlN grown with high purity Sc source reported by Casamento et al [106].  [4,18] are plotted as dashed curves. Experimentally measured values from sputtered [16,61,123] and MBE-grown [19,47,69] ScAlN films are depicted as scattered red and blue data points, respectively.

Bandgap and dielectric properties
The addition of Sc into nitride materials not only leads to lattice distortion but also results in direct-to-indirect bandgap crossover in Sc-III-N, which have been theoretically predicted and experimentally demonstrated (figures 8(a) and (b)) [4,47,61]. Theoretical calculations by Zhang et al indicate that the bandgap of Sc x Al 1−x N and Sc x Ga 1−x N stays direct up to x = 0.25 and x = 0.5, respectively [4]. In Sc x Al 1−x N, the Sc-3d states dominate at the conduction band minimum and lead to flat electron dispersion, and the bandgap decreases with increasing x. Conversely, the bandgap increases with increasing x for Sc x Ga 1−x N (up to x = 0.5), where Sc-3d states do not contribute to the conduction band minimum. However, due to the limited availability of materials with a comparable quality exhibiting pure phase over a wide composition range, deviations exist between theoretical and experimental data.
Preliminary results from optical absorption measurements showed that the bandgap of ScGaN decreases with increasing Sc content [65,78]. Based on the DFT calculations, Zhang et al pointed out the lower energy optical absorption onsets during measurements could stem from rock-salt-like crystal defects [4]. By growing ScGaN on different substrates, Tsui et al successfully demonstrated phase-pure wurtzite ScGaN films with their direct optical bandgap increased from 3.33 eV to 3.89 eV with increasing Sc content, and ascribed the decreasing bandgap to nanoscale lamellar inclusions of the zincblende phase, in agreement with theory [107].
Similarly, studies on the bandgap and optical properties of ScAlN are also significantly affected by the compositiondependent material quality and phase purity. The bandgap of rock-salt phase ScAlN grown on MgO substrate is found to be much smaller than wurtzite ScAlN [108,109]. Based on optical absorption results, Deng et al [110] suggested a bandgap of E g = 6.15-9.32x eV in wurtzite Sc x Al 1−x N with x < 0.2, which shows a significant deviation from the theoretically predicted values due to the limited material quality. By further improving the material quality, an empirical relationship between the bandgap and Sc content derived from Vegard's law for x < 0.25 was reported in sputtered and MBE-grown ScAlN films by Baeumler et al [61] and Wang et al [47], respectively ( figure 8(b)). Their experimental results agree well with early theoretical predictions [4]. For higher Sc concentrations, the bandgap is expected to become indirect, and the experimental data starts to slightly deviate from the theoretical values for a direct bandgap [47,61].
Boron-substituted AlN (BAlN) is generally considered as an ultrawide bandgap semiconductor [111,112], and the measured optical bandgap for low B alloyed BAlN exhibiting ferroelectricity is larger than 5.2 eV, which could be beneficial for reducing the dielectric loss in nitride ferroelectrics [113,114]. Hexagonal BN (hBN) has long been used as dielectrics for 2D heterostructures and devices, and the optical properties and bandgap have been extensively studied [115][116][117].
The band alignment of Sc x Al 1−x N/GaN heterojunctions with x = 0-0.25 has been investigated by first-principle DFT calculations. Both conduction band and valence band offset drop with increasing Sc content. Interestingly, the valence band offset approaches zero when Sc content approximates 25%, indicating a transition from type-I to type-II heterojunction [118], which was supported by recent experimental results [119]. Ambacher et al further predicted a 230% increase in the spontaneous polarization of ScAlN relative to AlN as a consequence of the deformed crystal structure and the iconicity of the Sc-N bond, which could be used to induce unprecedented interface charges up to 10 14 cm −2 in ScAlN/nitride heterostructures [18,120,121].
Beside the bandgap tuning with adding Sc, it is shown that the relative and high frequency dielectric constant of sputtered ScAlN also increases with Sc incorporation (figure 8(c)) [122][123][124]. Recent studies confirmed the increase in the permittivity of epitaxially grown ScAlN films relative to AlN, and pointed out the enhancement could come from an increase in the out-of-plane component of the ionic part of the dielectric constant [19,32,125]. The potentiality of ScAlN as highk dielectrics could be attractive not only to the nitride community, but also for silicon CMOS technology due to its promise to enhance transistor gate control. Note that although the dielectric constants of ScAlN are enhanced relative to AlN, they are still one order of magnitude lower than those of conventional piezoelectric materials like PbZr x Ti 1-x O 3 (PZT), advantageous for improving the electromechanical coupling in acoustic devices [122]. Detailed calculations on the piezoelectric, elastic, spontaneous polarization, and dielectric properties of ScAlN can be found in [18]. Likewise, the dielectric constants of ScGaN are also shown to increase with Sc incorporation [67].

Piezoelectricity
Modern wireless communication applications have recognized AlN as a well-established material for RF microelectromechanical systems with CMOS compatibility and high Curie temperature. However, compared with other piezoelectric materials, AlN exhibits relatively low piezoelectric coefficient and electromechanical coupling. Farrer et al [1] and Ranjan et al [2] predicted the phase transition from a nonpolar centrosymmetric hexagonal to a wurtzite polar structure when alloying Sc with wurtzite GaN or InN could lead to an increase in piezoelectricity. While predictions had been made on ScGaN and ScInN, the first experimental demonstration of enhanced piezoelectric response in Sc-alloyed III-nitride was based on ScAlN. In 2009, Akiyama et al reported a 400%-500% increase of the piezoelectric coefficient in sputtered ScAlN films [16], from 7 pC N −1 in AlN to 28 pC N −1 in ScAlN with 40% Sc [28]. Based on quantum mechanical calculations, Tasnadi et al showed that the giant piezoelectric moduli in ScAlN is intrinsic and comes from the softening of the C 33 elastic constant associated with the flattening of the energy landscape due to a phase competition between parent wurtzite and metastable hexagonal phase [3]. The piezoelectric response tends to increase with Sc content but decreases after around 50% Sc, as at higher Sc content, the material quality degenerates seriously due to phase separation, cubic inclusions, and elemental segregation [4,16,126,127]. Wingqvist et al [122] and Umeda et al [28] further confirmed the promise of superior electromechanical coupling and high sound velocity within wurtzite ScAlN films. Those findings, together with the modest dielectric constants and the high Curie temperature of ScAlN, have quickly driven the transition from AlN to Sc-alloyed III-nitrides in electromechanical devices for high frequency applications.
While the initial predictions of piezoelectricity enhancement had been made on ScGaN and ScInN, it was not until 2019 that enhancement in the piezoelectric response of ScGaN was reported, from 3.5 pC N −1 in GaN to 14 pC N −1 in ScGaN with 41% Sc [67]. Such an enhancement in ScInN, however, has not been reported yet. Recent experiments and calculations also showed that BAlN and yttrium (Y) alloyed III-nitrides could be ferroelectric [114,128]. Since all ferroelectrics are piezoelectric, it would be advantageous to explore the piezoelectric response limit in those emerging materials [129][130][131]. The typical piezoelectric coefficients of ScAlN and ScGaN are displayed in figure 9. The spontaneous polarization and piezoelectric constants of BAlN and BGaN have been theoretically studied in [132].
Additionally, the pyroelectric property is also enhanced in ScAlN, with up to a 25%-85% increase in effective pyroelectric coefficient compared to pure AlN [133,134]. A detailed study by Kurz et al indicates the additional external stress from the substrate, namely the clamping pyroelectric component, is the major contributor of effective pyroelectric coefficient, while the primary and secondary pyroelectric effects have  opposite signs and only contributes negligible pyroelectric charge, which are beneficial for stabilizing electroacoustic device operation [135]. The enhanced effective pyroelectric coefficients also arouse interest in Sc-III-N based thermal detectors [29].

Theoretical prediction of ferroelectricity in nitrides.
In crystal systems, the reported ferroelectric materials belong to the ten polar point groups ( figure 10(a)) [136,137]. Conventional wurtzite phase III-nitride alloys are piezoelectric and pyroelectric crystals, but not ferroelectric, as the orientation of their spontaneous polarization cannot be inverted by external electric field [41,99]. Considering all ferroelectric materials displays piezoelectric and pyroelectric behavior, although the converse is not ideally true, as outlined in figure 10(b), it is highly possible to realize ferroelectricity through lattice engineering in nitride semiconductors. Structurally, to achieve ferroelectricity in III-nitrides, their wurtzite phase (P6 3 mc) lattice should be able to switch between M-polar and N-polar ( figure 11(a)). This transition could happen through an intermediate hexagonal phase (P6 3 /mmc), which is the closest higher symmetry relative of the wurtzite structure. The main difference in these two structures is the internal parameter u, defined as the ratio of the length of the metal-nitrogen bond parallel to the c-axis and the lattice parameter c. For wurtzite structure, 0.375 < u < 0.5, and for the hexagonal phase, u = 0.5, indicating the metal and nitrogen atoms locating in different planes in the former, while sharing a common plane in the latter. Thus, the barrier energy for ferroelectric polarization switching in nitride materials can be determined as the energy difference between the equilibrium wurtzite phase and the metastable hexagonal phase. However, the calculated switching barrier for GaN was about 0.60 eV/formula unit, corresponding to an electric field, i.e. coercive field, of 23 MV cm −1 , which is significantly larger than the breakdown field of 5 MV cm −1 , suggesting ferroelectric switching cannot happen before dielectric breakdown [4]. To dramatically reduce the energy barrier, conventional III-nitrides have to be alloyed with other nitride compounds with favorable or metastable hexagonal phase. RE-nitrides, such as ScN, YN, lanthanum nitride etc, with metastable hexagonal phase [1,2] and hBN [115,138] are promising candidates.
Although ScN has a ground-state rock-salt phase [40], two metastable phases, i.e. tetrahedral coordinated wurtzite structure [42] and five-times coordinated hexagonal structure [1], have been predicted for ScN. This makes it possible for alloying conventional III-nitrides with ScN under wurtzite framework. Farrer and Bellaiche predicted that ScGaN and ScInN with low and even intermediate Sc content likely have a hexagonal phase, in which the disordered lattice can bridge the semiconductor and the ferroelectric classes [1]. In 2003, Ranjan et al predicted a nonpolar (hexagonal) to a polar (wurtzite) structure transition for ScN by applying compressive strain [2]. This phase transition process exhibits a switchable polarization, i.e. ferroelectricity. Unfortunately, the metastable hexagonal phase ScN has higher formation energy (0.316 eV −2 atoms) than the stable rock-salt phase [1], and the in-plane compressive strain adds more energy in the lattice, which make the material too unstable to be synthesized practically [2]. Therefore, they suggested to overcome this energetic problem by alloying ScN with GaN [2]. These theoretical works [1,2] suggested that ferroelectric polarization switching may likely be achieved in Sc-III-N alloys.
In 2010, Tasnadi et al attributed the origin of the abnormal increase of piezoelectric coefficient in ScAlN to a flattening of the energy landscape due to a competition between the wurtzite and hexagonal phases [3]. Shown in figure 11(b), with the addition of Sc in AlN, the wurtzite phase derived global energy minimum and the hexagonal phase derived energy saddle point was connected by a shallow region along the decreasing direction of c/a. Similar Gibbs free-energy flattening in perovskite ferroelectric materials has been discussed by Budimir et al in the framework of the Landau-Ginzburg-Devonshire theory [139]. Significantly, this flatten energy in ScAlN leads to a reduction of the energy barrier between the two stable polarization states of wurtzite ScAlN. Therefore, ferroelectric polarization switching in ScAlN is possible by applying an external electric field to overcome the phase transition barrier. In 2013, Zhang et al calculated the properties of Sc-III-N [4]. It was found that the c/a ratio showed a clear decreasing trend accompanied by an increase in u with increasing Sc content, indicating a closer distance between the planes containing metal atoms and N atoms in wurtzite phase Sc-III-N with higher Sc content, i.e. a lower polarization switching barrier. Therefore, Zhang et al suggested that ScGaN and ScAlN are ferroelectrics, in which the ferroelectric polarization switching barrier was directly related to u [4]. Even though those works provide a solid theoretical basis, ferroelectric polarization switching in ScAlN had not been experimentally demonstrated until 2019 by Fichtner et al [5]. After that, several groups performed more detailed analyses of the ferroelectric switching in Sc-III-N. In 2021, Wang et al reported a decreasing trend of ferroelectric switching barrier in ScAlN with the increase of both Sc content and temperature based on the Ginzburg-Landau theory (figure 11(c)) [140]. In 2022, Yazawa et al suggested that the strain sensitivity of switching barrier is one order of magnitude larger than that of the spontaneous polarization in wurtzite nitride ferroelectrics (figures 11(d) and (e)) [141]. Except for random Sc-III-N alloys, ferroelectric polarization switching properties in ScN/III-N superlattice has been also predicted by Noor-A-Alam et al [142], Sun et al [128], and Jiang et al [143] recently.
Similar flattening in potential-energy landscape as observed in Sc-III-N had been also predicted in wurtzite Y-III-N. In 2013, Tholander et al suggested that alloys exhibited a flatter potential-energy landscape and higher increase of piezoelectric coefficient when the parent components possessed close volume matching [129]. Therefore, they suggested that YInN, beyond ScAlN, as a promising piezoelectric material [129]. Very recently, Sun et al predicted a switching barrier as low as 0.042 eV for Y 0.5 In 0.5 N, corresponding to a coercive field of 1.61 MV cm −1 , which is significantly lower than that of Sc 0.5 In 0.5 N and Sc 0.5 Al 0.5 N [128]. It indicates that the wurtzite YInN is another promising ferroelectric material. In addition, adding other RE elements, such as lanthanide elements, into III-nitrides is possible to achieve ferroelectricity, in which further studies in both theory and experiment are needed.

Ferroelectric ScAlN.
While theoretical predictions have been made 20 years ago, the experimental demonstration of ferroelectricity in Sc-III-nitride was not possible until recently. Difficulty lies in reducing the switching energy barrier while preserving the dielectric breakdown strength of the material, which could be done by increasing the incorporation of Sc or strain engineering. In 2019, Fichtner et al shared the exciting news that ferroelectric switching is possible in sputtered polycrystalline ScAlN with Sc content up to 40% [5]. With increasing Sc incorporation, the coercive field decreases from ∼5 MV cm −1 to ∼2 MV cm −1 together with a slight remnant polarization reduction (figures 12(a) and (b)). The maximum extracted remnant polarization exceeds 100 µC cm −2 , which is consistent with theoretical predictions [98]. The polarity change has also been confirmed by piezoelectric response and wet etching experiments. In 2021 Wolff et al performed a detailed TEM study, and confirmed the switchable polarization for ScAlN from atomic scale (figures 12(c) and (d)) [60]. They investigated the electric field induced inversion domains, as well as the non-switched domains at the bottom electrode interface, providing critical information for understanding the domain evolution process during polarization switching in wurtzite materials. Subsequently, the impact of the nitrogen-to-argon gas ratio and target power ratios, plasma modes, sputter power, deposition temperature, crystal orientation, film stress, surface roughness, Sc content, and deposition rate on ferroelectric properties of ScAlN have been intensively investigated [62,64,[144][145][146][147][148][149][150][151][152][153].
However, sputtering deposition yields limited crystallinity with the presence of extensive defects and dislocations in the ScAlN films. Very recently, by using MBE, Wang et al have successfully demonstrated fully epitaxial, single crystalline, ferroelectric ScAlN on n-GaN over a Sc content range of 0.14-0.36 (figures 13(a) and (b)) [10]. The ferroelectric polarization switching in such epitaxially grown ScAlN films has been further confirmed using piezoresponse force microscopy measurements (figures 13(c) and (d)). Coercive field in the range of 3-6 MV cm −1 , and remnant polarization in the range of 120-150 µC cm −2 with a retention time beyond 10 5 s were obtained (figures 14(a) and (b)). Thanks to the improvement in crystal quality via epitaxial growth, the endurance strength was improved from 10 5 cycles, which is typical for sputtered ScAlN films [11], to more than 10 7 cycles (figure 14(c)) [32]. For ferroelectrics grown on a polar substrate, the polarization switching behavior could be altered by the polarity of the substrate [7]. Through growing nearly latticematched N-polar Sc 0.21 Al 0.79 N/GaN heterostructure on Npolar GaN, Wang et al reported a highly uniform coercive field (∼4.6 MV cm −1 ) and remnant polarization (∼90 µC cm −2 ) across the whole wafer [125], comparable to those obtained on M-polar substrates [10]. In conventional III-nitride heterostructures, the existence of dislocation-related leakage paths has been shown to greatly affect the breakdown strength of the material. By growing ScAlN epilayers on bulk GaN, GaN on sapphire, and GaN on Si substrates with dislocation densities ranging from 10 7 to 10 10 cm −2 , Wang et al  reported that the existence of high densities of dislocations does not smear out the presence of ferroelectricity, but rather results in higher leakage current, reduced breakdown strength, and overestimated remnant polarization values [15]. The epitaxial nature in MBE growth could enable integration of highperformance ferroelectric functionality with well-established semiconductor platforms for a broad range of electronic, optoelectronic, and photonic device applications.
Following these initial demonstrations, the thickness scaling down effect in nitride ferroelectric semiconductors has been studied. Fichtner et al investigated the thickness scaling and imprint behavior of ferroelectric ScAlN and demonstrated that ferroelectricity is possible for ScAlN below 30 nm thickness while the coercive fields were found to be largely independent of thickness down to 27 nm [154]. Most recently, their group reported ferroelectric switching in 10 nm Sc 0.28 Al 0.72 N films, in which the whole ScAl-N/Pt heterostructures were in situ sputtered onto GaN/sapphire substrates [155]. Wang et al successfully deposited 100 nm-thick ferroelectric ScAlN layer directly on Pt without amorphous region or dead layer [144], and further showed ferroelectric switching in 20 nm-thick ScAlN films with a Sc content of 0.32 and 0.36 [156] using an alternative measuring setup [30]. Yasuoka et al reported the effect of sputtering deposition conditions on the ferroelectric behaviors of ScAlN, and achieved ferroelectricity in 9 nm-thick ScAlN film from scanning nonlinear dielectric microscopy measurements [62,157], which was further investigated by positive up negative down (PUND) characterizations showing a remnant polarization beyond 100 µC cm −2 [157,158]. Tsai et al also observed ferroelectric-type hysteresis behavior in sputtered Sc 0.28 Al 0.72 N films with a thickness down to 20 nm using capacitance measurements, while the remnant polarization showed a gradual degradation when the thickness is less than 35 nm [159]. Ryoo et al reported ferroelectricity in sputtered ScAlN films with a thickness down to 20 nm by using optimized process conditions [160]. Most recently, we also demonstrated ferroelectric polarization switching in MBE-grown ScAlN films with a thickness of ∼5 nm on n-GaN/sapphire and Mo(011) substrates [161,162]. The experimentally reported thickness dependent remnant polarization and coercive field for ferroelectric ScAlN films by different research groups are summarized in figures 15(a) and (b), respectively. Although the ferroelectric polarization switching has been observed experimentally, more studies on the effect of interfacial bonding, dead layer, stress, native oxide layer, etc Nitride ferroelectrics have a high Curie temperature nature compared to conventional ferroelectrics [163]. Temperature dependent measurements indicated that the coercive field can be significantly reduced with elevating temperature, which results in three times frequency tunability in ScAlN based acoustic devices compared to AlN and potentially enable low voltage operation [12,164,165]. A more detailed study was conducted by Zhu et al through comparing the temperature dependence of ferroelectric switching in ScAlN, AlBN, and AlN [166]. Almost 50% drop of coercive field was observed for ScAlN at 300 • C ( figure 15(c)), and based on thermal activation model a distribution of pinning sites or nucleation barriers regulated switching process was suggested. To test the high temperature limit of wurtzite ferroelectric ScAlN, the wurtzite structure of ScAlN has been studied with high temperature annealing cycles, showing a ferroelectricto-paraelectric transition temperature surpassing 1100 • C, and excellent retention over 10 5 s even at high temperatures [11,163,167,168].
Currently, leakage current remains a serious issue in ferroelectric ScAlN films. Tsai et al examined the large leakage current in sputtered ScAlN films and ascribed the leakage current to the formation of nitrogen vacancies at the metal electrode interface [169], which could be aggravated in thinner films or after excessive cycling and could cause breakdown via Joule heating at local conductive filaments [159,170]. By using a compositionally modulated periodic multilayer ScAlN structure, Zheng et al observed an increase in breakdown electric field, which was understood as a blocking or slowing effect of the electrical tree, i.e. leakage path, by the multilayer interfaces [171].

Ferroelectric ScGaN
As a sister-material of ScAlN, theoretical studies predicted that the incorporation of Sc in GaN could also enhance the piezoelectric response and reduce the polarization switching barrier in the manner similar to that of ScAlN [67,172]. It has been demonstrated that the c/a ratio of ScGaN tends to decrease with increasing Sc content in concert with enhanced piezoelectric moduli [67,78]. The first demonstration of ferroelectric ScGaN was reported by Wang et al in 2021, which was grown on n-GaN using MBE (figures 16(a) and (b)) [83]. A wake-up process was revealed during the first several measurement cycles, and after sufficient poling, coercive fields in the range of 2.0-3.0 MV cm −1 , and large, retainable remnant polarization in the range of 60-160 µC cm −2 were measured for ScGaN thin films with Sc contents in the range of 0.31-0.41. Analogous to ScAlN, the coercive field was found to decrease with increasing Sc incorporation. Later, Uehara et al also confirmed the ferroelectricity in sputtered ScGaN (figures 16(c) and (d)), and reported a remnant polarization exceeding 120 µC cm −2 and coercive field of 3.6 MV cm −1 for Sc contents in the range of 0.35-0.44 [68,173]. By comparing the Sc content dependent coercive field in ScAlN and ScGaN, they pointed out the Sc content governed the coercive field value, suggesting that the polarization switching started in the wurtzite unit cells containing Sc, which is consistent with a recent theoretical prediction by Krishnamoorthy et al [174].

Ferroelectric BAlN.
Similarly, BN is stable in its hexagonal (hBN) and cubic form, and alloying III-nitrides with B is also possible to flattening the energy landscape. In 2021, Hayden et al reported the demonstration of ferroelectricity in B-substituted AlN thin films and the coercive field was found to decrease with increasing B concentration (figures 16(e) and (f)) [114]. BAlN thin films with B contents in the range of 0.02-0.15 displayed ferroelectric polarization switching with a remnant polarization exceeding 125 µC cm −2 and a coercive field larger than 5 MV cm −1 [114]. In their following studies, Zhu et al reported a weak-up process in the ferroelectric B 0.07 Al 0.93 N films [175]. During this weak-up process, the remnant polarization gradually increased from 0 to >100 µC cm −2 , and it was also accompanied with an increase of the density of domain walls [175].

Ferroelectric hBN moiré lattice.
Very recently, stable ferroelectric order has been observed at the interface between two hBN flakes, which are stacked together in a metastable non-centrosymmetric parallel orientation via a van der Waals force [176][177][178][179][180]. The origin of spontaneous net electric polarization in bilayer hBN is illustrated in figure 17(a) [177,178]. Bilayer hBN has six different high-symmetry stacking configurations, which are divided into two groups with 'parallel' and 'antiparallel' orientations. Typically, hBN crystal naturally grows in the energetically favorable antiparallel (AA ′ ) configuration, i.e. the nitrogen (boron) atoms of one layer are fully overlapped with the boron (nitrogen) atoms of the adjacent layer, resulting in zero net electric polarization. In addition, the AA ′ configuration is more favorable than the AB ′ and BA ′ configurations in a relaxed system with antiparallel orientation [178]. While in the parallel orientation, fully overlapped configuration (AA) is unstable. Instead, the half-overlapped AB and BA stacking with equivalent lattice structures are the energetically favorable metastable configurations. Due to the asymmetric charge distribution, spontaneous polarization with opposite direction occurs in AB and BA configurations. This suggests that non-zero net spontaneous polarization can be constructed via stacking engineering in hBN. To explore this, researchers stamped two exfoliated hBN flakes together with a minute twist angle [176][177][178][179][180]. The small twist leads to continuous interlayer translation and moiré lattice formation owing to the lattice periodicity (figure 17(a) middle). With the structural relaxation process, commensurate large domains with AB and BA stacking configurations appear at adjacent regions and are separated by sharp incommensurate AB/AA/BA domain walls [177]. Figure 17(b) presents a Kelvin probe force microscope (KPFM) image measured from stacked hBN flakes, showing oppositely polarized domains of AB/BA stacking (black and white), ranging in area between ∼0.01 and 1 µm 2 and separated by sharp domain walls [177]. Furthermore, redistribution of domain walls was observed by applying negative/positive bias, suggesting these domains have ferroelectric-like switchable polarization [177]. The ferroelectricity of hBN persists to sub-micrometer scale domains, which greatly extends the application landscape of nitride ferroelectrics and paves the way for potential ultrathin memory devices [176,179,180].
Extensive studies have provided unambiguous evidence that nitride ferroelectric material system is a thin film ferroelectric with large remnant polarization, high coercive field, and exceptional temperature stability and durability compared to common ferroelectric material classes ( figure 18). The demonstration of ferroelectricity in nitride semiconductors could enable a broad range of emerging applications with combined functionality in ferroelectric, electronic, optoelectronic, photovoltaic (PV), photonic, and quantum devices and systems.
Aside from ferroelectric properties, Lei et al reported roomtemperature ferromagnetic behavior in Sc-doped AlN sixfoldsymmetrical hierarchical nanostructures, where the ferromagnetism is associated with Sc 3+ introduced Al vacancies [181]. Through chemical-doping engineering, Liu et al demonstrated the coexistence of ferroelectricity and ferromagnetism in Nidoped Sc 0.3 Al 0.7 N thin films, opening the avenue for nitride multiferroic materials [182].

Nonlinear optical properties
It is well known that the propagation direction, spatial profile, and transmission of an optical field through a material can be modified by the light-matter interaction, which is determined by the electric polarization induced by the applied light [183,184]. In the case of linear optics, the polarization depends linearly on the electric field. However, the light induced polarization can be expressed as a power series in the electric field in the nonlinear optics, in which χ (1) is the linear susceptibility, while χ (2) and χ (3) are the second-and third-order nonlinear optical susceptibilities of the material, respectively [183,185]. Recently, nonlinear optical materials have attracted tremendous interest in a wide range of applications, such as photonics, nano-photonics, bio-photonics, and quantum photonics [186,187]. Interestingly, most ferroelectrics also possess a large nonlinear response with respect to the electromagnetic radiation in the optical range [188][189][190]. As discussed above, nitride ferroelectrics have a wurtzite crystal structure, which is a noncentrosymmetric structure [4], enabling both the χ (2) and χ (3) optical nonlinearity. As seen in the case of piezoelectric enhancement in ScAlN, it is expected that the optical nonlinearities can also be boosted. Recently, Yoshioka et al characterized the second-order nonlinearities of sputtered ScAlN films with a Sc content up to 36%. It was found that χ (2) increased with the increase of Sc content [20]. The χ (2) component d 33 was measured to be 62.3 ± 5.6 pm V −1 for Sc 0.36 Al 0.64 N, which shows more than 12 times improvement than pure AlN (5.1 ± 0.4 pm V −1 ) and is 2 times higher than the extensively studied LiNbO 3 (25.2 pm V −1 ) [20,191]. This significant enhancement is likely directly related to the lattice distortion in wurtzite ScAlN [3,16,18]. Consequently, such intriguing properties make it a very promising material in nonlinear optical device applications.

HEMTs
Addition of Sc also introduces larger spontaneous polarization into ScAlN besides the abnormal enhancement of piezoelectric response [4,18]. For example, the spontaneous polarization of Sc 0.18 Al 0.82 N (∼0.12 C m −2 ), which is lattice matched to GaN, is about 1.3 and 3.5 times of that of AlN (∼0.09 C m −2 ) and GaN (0.034 C m −2 ), respectively [18]. In III-nitride material system, a 2DEG or 2DHG is formed when a net polarization charge is presented at the heterointerface, wherein the charge density in the channel is controlled by the magnitude of the net polarization [192,193]. AlGaN/GaN, AlInN/GaN, and InGaN/InN heterostructures based HEMTs have been extensively studied as state-of-the-art high power and high frequency transistors [194][195][196][197][198][199]. Therefore, analogous to AlGaN and AlInN, ScAlN is a promising barrier material in GaN HEMTs. Figure 19(a) shows the simulated sheet charge density for GaN HEMTs with different barrier materials [200]. Due to the lattice mismatch induced piezoelectric polarization, AlN/GaN has the highest sheet charge density. Unfortunately, the critical thickness for AlN grown on GaN is only 5 nm due to the significant tensile strain. In contrast, Sc 0.18 Al 0.82 N not only exhibits a higher sheet concentration, but also is lattice matched to GaN, thereby eliminating the critical thickness limitation.
In 2017, Hardy et al demonstrated the 2DEG in a MBEgrown Sc 0.14 Al 0.86 N/GaN heterostructure with GaN/AlN interlayer [45]. A sheet charge density of 3.14 × 10 13 cm −2 and a Hall mobility of 910 cm 2 V −1 s was measured. In 2019, Frei et al reported a sheet carrier density of 5.26 × 10 13 cm −2 with Hall mobility of 553 cm 2 V s −1 in a MBE-grown Sc 0.20 Al 0.80 N/GaN heterostructure by optimizing growth conditions [46]. Most recently, utilizing MBE, Casamento et al found that a ∼2 nm AlN interlayer boosts the room temperature mobility by more than five times from 300 cm 2 V s −1 for a ScAlN/GaN heterojunction to 1573 cm 2 V s −1 for a ScAlN/AlN/GaN heterojunction [201]. The temperature dependent Hall mobility and sheet carrier density of these two kinds of heterojunctions are compared in figures 19(b) and (c), respectively. In 2019, the presence of 2DEG was reported in MOCVD-grown ScAlN/GaN HEMT structures [51], in which the electron mobility was further improved by reducing the concentration of carbon and oxygen impurity through elevating growth temperature [52]. Aside from these heavily studied M-polar ScAlN/GaN HEMTs, Wang et al demonstrated N-polar ScAlN/GaN HEMTs on sapphire substrates in 2021 [24]. The sheet electron concentration varied from 3.86 × 10 13 to 4.88 × 10 13 cm −2 with increasing ScAlN barrier thickness ( figure 19(d)). In an N-polar GaN HEMT structure with a Sc 0.18 Al 0.82 N barrier thickness of 15 nm, a sheet electron concentration of 4.1 × 10 13 cm −2 and a Hall mobility of 564 cm 2 V s −1 was measured, and the corresponding sheet resistance was as low as 271 Ω sq −1 .
With the demonstration of 2DEG in ScAlN/GaN heterostructures, HEMT devices have been also fabricated and characterized shortly. In 2019, Green et al reported the device performance of MBE-grown ScAlN/GaN HEMTs [22]. Employing a highly doped n + -GaN regrowth approach, they achieved an ohmic contact resistance as low as 0.1 Ω mm. Figures 19(e) and (f) display the output curves and DC transfer characteristics of the ScAlN/GaN HEMT devices, showing a maximum current density of 2.4 A mm −1 and a peak transconductance of 0.67 S mm −1 [22]. In 2020, Green et al further compared the RF power performance of ScAlN/GaN and ScAlGaN/GaN HEMT devices [23]. The ScAlGaN/GaN HEMTs delivered an output power of 5.77 W mm −1 (V D = 20 V) and a power-added efficiency of 47% (V D = 15 V) with 102 GHz maximum oscillation frequency. Very recently, the improvement of material quality as well as 2DEG mobility in ScAlGaN/GaN compared to ScAlN/GaN has been reported [69], which maybe one of the reasons for the performance improvement in ScAl-GaN/GaN HEMT devices. Most recently, Krause et al reported an output power of 8.4 W mm −1 and a power-added efficiency of 48% in MOCVD-grown ScAlN/GaN HEMT devices [202].
Achieving transistors with higher speed at lower power dissipation has been the undergoing goal of Moore's law. However, due to the Boltzmann distribution of electrons, there is a fundamental barrier to further reducing the power dissipation in conventional electronic devices, i.e. the minimum limitation of subthreshold swing (SS ⩾ 60 mV dec −1 ), often termed as Boltzmann's Tyranny [203][204][205][206]. In 2008, Salahuddin and Datta suggested that the negative capacitance (NC) in a ferroelectric capacitor can be used as a solid solution for overcoming the Boltzmann's Tyranny in electronic devices [203]. In the past years, NC has been experimentally measured in ferroelectric PZT [207], PbTiO 3 (PTO) [208], Hf 0.5 Zr 0.5 O 2 (HZO) [209], etc [210]. Recently, field effect transistors (FETs) with SS lower than 60 mV dec −1 have been experimentally demonstrated by using ferroelectric dielectric layers [211]. Even though the physical picture and design rule for NC-FETs are still under debate [212], this field has attracted significant interest. The appearance of emerging nitride ferroelectrics, especially its epitaxial integration ability with conventional III-nitride transistors and its advantages in HEMTs [10,24,83], offers new opportunities for the design and fabrication of NC-FET devices with lower power dissipation.

Ferroelectric memristors
Due to the increasing demand in storing, retrieving, and computing a large volume of information, the interest in ferroelectrics is refueled due to their promising applications in energy efficient memory, neuromorphic and in-memory computing, and edge intelligence. The electrically reconfigurable, nonvolatile and large spontaneous polarization of nitride ferroelectrics can be used to non-volatilely modulate the barrier height and width in a metal-ferroelectric-metal capacitor or a metal-ferroelectric-semiconductor heterostructure to build resistive memristors and programmable homojunctions. Liu et al demonstrated a large difference in leakage current as a function of ferroelectric polarization in a Pt/ScAlN/Pt capacitor with 20 nm-thick ScAlN, which has been attributed to an electron hopping process associated with the trap charge states [31]. A MoS 2 /ScAlN back-barrier ferroelectric transistor was also demonstrated by Liu et al (figures 20(a)-(c)) [13]. Thanks to the large coercive field and remnant polarization of ScAlN, the ScAlN/2D channel ferroelectric FET showed wide memory window and high ON/OFF ratio, which is promising for embedded memory and memory-based computing architectures [13]. Most recently, Liu et al further reported a transistor-free compute-in-memory architecture that permits storage, search, and neural network operations on ScAlN ferroelectric diodes [213]. Based on DFT calculations and electrical measurements, Liu et al pointed out N vacancies are more likely to exist and dominate the leakage current in ScAlN [214]. It has been demonstrated that in ferroelectrics vacancy migration could occur and contribute to the resistive switching or potential modulation [215]. Therefore, polarization coupled measurements were proposed to distinguish and examine the contribution from different mechanisms [216]. In 2022, Wang et al reported the first epitaxial ferroelectric ScAlN/GaN heterostructure memory (figures 20(d)-(f)) [32]. The structure exhibited robust high and low resistance operation with ON/OFF ratio in the range of 60−210, retention time of over 3 × 10 6 s, and bipolar cycling of over 10 4 times. By conducting detailed polarization coupled resistance measurements and tuning the doping level of the GaN electrode, the resistive switching behavior is directly linked to polarization switching, and successfully ascribed to the ferroelectric field effect which causes charge reconstruction at the ferroelectric/semiconductor interface. Remarkably, it was found that the heterostructure maintained a rectifying ratio larger than 10 even at 400 • C, suggesting ScAlN based memristors could be potentially high temperature memory for harsh environments.
Most recently, by growing GaN/ScAlN/Mo heterostructures, we were able to demonstrate a semiconductor/ferroelectric/Metal memristor with the opposite ON/OFF operation compared with above Metal/ScAlN/GaN configuration, further confirming the ferroelectric tuning of the interface barrier height [217]. This ferroelectric heterostructure also shows capability of emulating the spike-time-dependent plasticity in a biological synapse. Furthermore, we also achieved large ON/OFF ratios towards analog in-memory computing using an ultrathin ScAlN/GaN ferroelectric heterostructure [161]. These results suggest that ScAlN-based ferroelectric/III-nitride heterostructures can be competitive candidates for next-generation memristors and all nitridebased monolithic integrated logic circuits for power-efficient applications and harsh environments.
One possible drawback with ScAlN related memories is the relatively high coercive field, which results in large operation voltages and potentially higher power consumption. While most of the reported ScAlN-based memories utilized relatively thick ScAlN layers, significant scaling down would be of great interest. Besides, the modest dielectric constant and giant remnant polarization would raise issues with depolarization field. We highlight the compatibility and heterogeneous integration of nitride ferroelectrics with both GaN and silicon technologies-the two most widely developed semiconductor technologies, which could deliver computation and logic functionality in beyond-silicon technology, further extending the landscape of data storage and computing.
Compared with ScAlN, ScGaN possesses relatively smaller bandgap and lower coercive field, which potentially could be used as ferroelectric semiconductor for novel memory configurations [218]. Similarly, BAlN could benefit from its large bandgap to tailor leakage in device applications. Considering the approximation in the crystal structure and lattice constants between the reported nitride ferroelectrics, one would expect a wealth of innovative structure designs based on the polarization and band profile engineering of nitrides for advanced device applications.

Ferroelectric PVs
Propelled by the escalating demand for sustainable and energy efficient technology, ferroelectric self-powered PV devices have drawn considerable attention, particularly for solar energy scavenging and photodetection [219][220][221]. Currently oxide perovskite ferroelectrics, such as BaTiO 3 , BiFeO 3 , and Pb(Zr,Ti)O 3 , are being used for those applications [222,223]. These materials are excellent photosensitive semiconductors but are limited by high density of defects, large recombination loss, large leakage/dark currents, as well as scalability issues [224]. Additionally, their narrow bandgap is unsuitable for solar blind UV photodetection. Epitaxially grown ferroelectric ScAlN, with its direct wide bandgap and large inherent polarization field, shows great promise for PV applications. Most recently, Mondal et al reported high-performance ferroelectric self-powered deep UV photodetectors based on epitaxially grown ScAlN thin films (figure 21) [225]. The responsivity at 193 nm illumination reached up to a maximum of 15 mA W −1 with a detectivity of 1.2 × 10 11 Jones at an illumination intensity of 0.12 mW cm −2 . Furthermore, the photodetectors exhibit wake-up free and reconfigurable photo-response, and fast and stable switching response time (<0.06 s) with excellent rejection to UVA and visible illumination.

Resonators and filters
The rapid advancements in communication technology for 5G and beyond have been fueled by the exponential growth of wireless users and Internet of things [226]. Acoustic resonator and filter technologies have shown to be critical to meet the ever-escalating demand of higher data rates, multi band transceivers. Sc 0.4 Al 0.6 N has shown a 400%-500% improvement in the piezoelectric coefficient, d 33 , which further translates to an increase of 30% in the electromechanical coupling coefficient, κ 2 [16,227,228]. This enhancement makes ScAlN to be a promising candidate for the realization of Lamb wave resonators [229]. Utilizing the high d 33 [232]. Following on that, Park et al demonstrated composite thin-film BAW resonators based on a sputtered ScAlN thin film [233]. They reported a high κ 2 value of 11.7% at the third-order thickness-extensional resonant frequency of 2.4 GHz, with a high κ 2 × Q max figure-of-merit (FoM) of 84 which is two times higher than the FoM reported by Wang et al [232]. Most recently, Zheng et al demonstrated the first SAW resonators using MBE-grown ferroelectric ScAlN/GaN heterostructure, in which high-Q factor and high-order Sezawa modes were probed, exhibiting a Bode Q factor of 571 at 4.6 GHz [234]. Additionally, the piezoelectric response changes together with the polarization switching in ferroelectrics. The combination of piezoelectricity and ferroelectricity enables new functionalities. The dielectric and piezoelectric properties of ferroelectric ScAlN thin films have been studied by Pirro et al, and the capacitance variation for different Sc contents were shown capable of canceling out inductance variations of up to 10% in impedance matching networks [235]. Wang et al showed that the sign of e 33 piezoelectric coefficient in ScAlN can be switched together with a 20 times change in the series resistance of the capacitors, and polarizationdependent operating states can be achieved in ScAlN based acoustic wave filters and resonators [27,232,236]. These demonstrations reaffirm the potential of ScAlN for the realization of highly integrated filters for multi-band 5G and beyond systems.

Nonlinear photonic devices
Strong optical nonlinearities provide various benefits in onchip quantum photonic devices, and plenty of applications have been established based on the AlN platform, such as frequency comb [237], electro-optic modulator [238], optical parametric oscillator [239], and ultrahigh Q resonators [240,241]. Due to the drastic enhancement of optical nonlinearities in ScAlN, employing ScAlN in the nonlinear optical devices can potentially improve conversion efficiencies, modulations, and other performances compared to AlN. Recently, Zhu et al reported ScAlN photonic circuits on Si, in which a waveguide propagation loss of 9 ± 2 dB cm −1 and microring Q factor of 1.4 × 10 4 was realized [242]. Zhang et al further simulated the performance of ScAlN-based waveguides with different geometric parameters and Sc contents at the near-infrared wavelengths [243]. It was found that optical mode can be well confined and single-mode operation can be achieved in the ScAlN waveguides by appropriately choosing the key design parameters. Yet the loss remains higher than AlN and other low-loss materials [240,241]. In addition, the etching process for ScAlN is still a challenging issue due to the formation of ScCl 3 , which has very low volatility during the dry etching step [244]. Therefore, more investigations are needed to achieve low-loss devices on the ScAlN platform to better exploit its high nonlinearities. In this context, we have recently demonstrated, for the first time, fully epitaxial ferroelectric ScAlN waveguides and microring resonators [245].

Summary and prospects
In previous sections, we have reviewed and presented some of the significant developments and achievements in the field of nitride ferroelectrics. To conclude, we survey some of the promising directions and potential challenges for the future developments of nitride ferroelectric materials and devices.
From the material point of view, the recent demonstration of single crystalline nitride ferroelectric semiconductors has shown the feasibility of developing a broad range of new ferroelectrics by incorporating other RE elements, from Y to La to Lutetium (Lu), in III-nitride semiconductors, offering unprecedentedly rich opportunities for strain, alloy, quantum, and entropy engineering. Besides, by alloying with ferromagneticactive elements via chemical-doping engineering, multiferroic nitrides could be constructed for subsequent research and related applications. The interfacial ferroelectricity by stacking engineering of bilayer 2D materials will welcome more material systems.
From the device point of view, an intuitive approach is to combine nitride ferroelectrics with or to use ferroelectrics to replace part of the materials in the state-of-the-art IIInitride devices like HEMTs, light emitting diodes (LEDs), laser diodes, photodetectors, waveguides, resonators, etc, to enhance device performance and tunability, and to add programmability, data storage, and computing capabilities. The fully epitaxial manner of ferroelectric Sc-III-N with other nitride materials further enable polarization engineering beyond conventional III-nitrides and allow for novel memory device designs including ferroic junctions and transistors taking advantage of the large polarization discontinuity and widely tunable bandgap in nitride semiconductors. Recently, the NC effect in ferroelectric materials has attracted significant attention. It would also be beneficial to explore the applications of nitride ferroelectrics in steep-slope transistors especially in the nitride material systems.
The main challenge for ferroelectric applications of ScAl(Ga)N is the large leakage current, which could originate from defects especially N vacancies, impurities, dislocations, and non-uniform domains. Identifying the electric leakage path in nitride ferroelectrics and developing strategies to suppress the leakage current are essential for future applications. Theoretical calculations, in combination with experimental studies, are needed to identify the defects type and how they impact the electronic, electromechanical, and optical properties. There is also a need to overcome the thickness scaling challenge of ScAlN down to a few nanometers, or possibly down to the monolayer limit, where detailed experimenttheory coupled investigations on the microstructures, switching dynamics, and energy relaxation processes at ferroelectric/semiconductor and ferroelectric/metal interfaces will be of significant interest. As this new class of ferroelectrics is still in its infancy, studies on the fundamental properties, like fatigue, imprint, and possible dead layer formation will be critical.
The discovery of ferroelectricity in Sc-alloyed III-nitrides could be a textbook example of designing multi-functionality through alloying dissimilar materials and a successful demonstration of theory-to-experiment methodology. As the discovery of ferroelectricity in hafnia-based materials in 2011 renewed the interest in ferroelectric memory devices, the discovery of nitride ferroelectrics, with their appealing properties inherited from III-nitride semiconductors, could fuel a new wave of intensive research on ferroelectric devices in a much wider range of applications including memory, power electronics, piezoelectronics, optoelectronics, and quantum photonics. With all those promises and potentialities, we are foreseeing a diversity of intriguing materials and devices to come.

Data availability statement
The data that support the findings of this study are available upon reasonable request from the authors.