Polarized neutron reflectometry is used to probe the in-plane projection of the net-magnetization vector M of polycrystalline Fe films exchange coupled to twinned (110) MnF 2 or FeF 2 antiferromagnetic (AF) layers. The magnetization reversal mechanism depends upon the orientation of the cooling field with respect to the twinned microstructure of the AF, and whether the applied field is increased to (or decreased from) a positive saturating field; i.e., the magnetization reversal is asymmetric. The reversal of the sample magnetization from one saturated state to the other occurs via either domain wall motion or magnetization rotation on opposite sides of the same hysteresis loop.

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MnF 2 ͞Fe bilayers exhibit asymmetric magnetization reversal that occurs by coherent rotation on one side of the loop and by nucleation and propagation of domain walls on the other side of the loop. Here, we show by polarized neutron reflectometry, magnetization, and magnetotransport measurements that for samples with good crystalline "quality" the rotation is a two-stage process, due to coherent rotation to a stable state perpendicular to the cooling field direction. The result is remarkably asymmetrically shaped hysteresis loops. DOI: 10.1103/PhysRevLett.86.4394 PACS numbers: 75.70.Cn, 75.30.Gw Despite recently renewed interest in the phenomenon of exchange anisotropy [1] at the interface between antiferromagnets (AF) and ferromagnets (F), a full understanding is elusive [2]. A complete description of exchange anisotropy would incorporate an explanation of the mechanism for exchange biasing [2][3][4][5][6][7][8], the effects of interface disorder [9][10][11], the relationship between the exchange bias ͑H E ͒ and the coercivity ͑H C ͒ [6,12,13], as well as an understanding of the magnetization reversal mechanisms [14][15][16]. In an attempt to address these issues, we have focused on a model thin film system TMF 2 ͞Fe (TM transition metal), with which we have been able to elucidate perpendicular coupling [17], effects of compensation [18] and interfacial disorder [9,10], and the relationship between H E and H C [12].A recent thrust has been to understand the mechanisms by which the magnetization reverses in such systems [15]. This is a fundamental point, clearly related to the behavior of H C (zero moment half-width of the hysteresis loop), which is considered important [6]. Although there are few more fundamental questions about the nature of a hysteresis loop than the issue of the magnetization reversal mechanism, it has largely escaped investigation in exchange biased systems. Recent papers [14,15] have contributed to the realization that the magnetization reversal is intrinsically asymmetric under certain conditions. This naturally explains the often observed [1,[18][19][20] [22]. In summary, it is clear that reversal asymmetry in exchange biased systems is a general phenomenon requiring immediate attention.In the MnF 2 ͞Fe and FeF 2 ͞Fe systems, we have shown that for certain cooling field orientations the magnetization reverses by coherent rotation on the high field side of the loop (i.e., where the coercive field is jH E j 1 H C ) and by domain wall nucleation and propagation on the low field side of the loop (i.e., where the coercive field is H C 2 jH E j), a very clear asymmetry [15]. The magnetization hysteresis loops show very slight shape asymmetry, which must be associated with this effect [23]. The asymmetry is easily understood on consideration of the twinned nature of the AF films (Fig. 1). When the cooling field bisects the twin anisotropy axes (as shown), the frustration of the coupling between the individual AF twins and the ferromagnetic overlayer leads to an effective "45 ± coupling," r...

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Polycrystalline CrN/TiN superlattice films were deposited on M1 tool steel using unbalanced reactive magnetron sputtering with opposed cathodes. The Cr and Ti targets were sputtered in Ar–N2 mixtures with partial pressure control of the N2. As the N2 partial pressure was increased from 0.1 to 1.1 mTorr, TiNx films went from stoichiometric B1-cubic TiN to slightly overstoichiometric TiN, while CrNx films went from cubic Cr–N solid solutions to hexagonal Cr2N to B1-cubic CrN. Since the N2 partial pressure required to form stoichiometric CrN was ≈10 times that required to form stoichiometric TiN, nitrogen was inlet at the Cr target position to maximize the difference in N2 partial pressures. Two series of CrN/TiN superlattices, with TiN fractions of 0.4 and 0.6, were deposited with periods ranging from 2 to 60 nm. X-ray diffraction showed a very strong (111) texture with first-order satellite peaks around the (111) Bragg peak. Kinematical diffraction simulations of the superlattice x-ray patterns indicated a strong composition modulation and a significant fluctuation in d-spacing that was related to ion bombardment defects. Cross-sectional transmission electron microscope images showed a columnar film structure with well-defined superlattice layers. Nanoindentation of 2-μm-thick CrN/TiN samples showed a maximum hardness of 35 GPa at a period of 2.3 nm, compared to 25 GPa for TiN and 14 GPa for CrN films. The maximum superlattice hardness was thus ≈75% larger than the rule-of-mixtures value. The hardness enhancement mechanisms are discussed.

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