Z alloys, the compositions of which show promise for manufacturing high-coercivity permanent magnets, is reported. Foundations arising from the quantitative analysis of alloy microstructures as the first, coarse, level of heterogeneity are considered. The structure of the alloys, in optical resolutions, is shown to be characterized by three structural phase components, which are denoted as A, B, and C and based on the 1:5, 2:17, and 2:7 phases, respectively. As the chemical composition of alloys changes monotonically, the quantitative relationships of the components A, B, and C vary over wide ranges. In this case, the hysteretic properties of the Z alloys in the high-coercivity state are strictly controlled by the volume fractions of the A and B structural components. Based on quantitative relationships of the A, B, and C structural components for the Z alloys with R = Gd or Sm, sketches of quasi-ternary sections of the -R-Zr phase diagrams at temperatures of 1160–1190 °C and isopleths for the 2:17–2:7 phase composition range of the –Sm–Zr system were constructed.An original vision for the structural formation of  Z alloys, which were developed and commercialized a while ago [c” axis, which is common for the whole anisotropic massive of samples [Permanent magnets based on the hile ago ,2, are chile ago ,4, have hile ago ,4,5,6 mehile ago . For exahile ago  assume t samples ,8. WithiZ alloys. This hinders the progress in the development and improvement of both new compositions of alloys for permanent magnets and manufacturing processes for them.It should be recognized that currently there is no objective holistic view of the formation of high-coercivity structure of the Z alloys in the composition ranges, which are of importance for the manufacturing of permanent magnets because of the functionality of their phase structure. The magnetic hardness of the Z alloy for permanent magnets is ensured by its precipitation hardening in the course of complex heat treatment. The detailing is performed based on the generalizing of both original results, which are not given in widely available periodicals and the known literature data. The study is also aimed at concretizing the phase transformations of the Sm–Zr–Co–Cu–Fe system in accordance with our understanding of the occurred processes and phenomena.The aim of the present study was to detail the interrelations of the chemical and phase compositions of the Z alloys, there are investigations related to Ray’s metallurgical model [The starting point is the matrix, which is the disordered single-phase 2:17R precursor; the decomposition of the phase during isothermal tempering leads to the formation of the cellular-lamellar structure of phases. The formed phases are modified at the expense of interphase interdiffusion during stepped aging.2 dumbbell sites when entering the Sm217 matrix.Zirconium with a vacancy occupies CoThe composition of the alloy for manufacturing the permanent magnets should be maximally close to the 2:17 stoichiometry taking into account Sm and Zr losses for the oxidation and reactions with other nonmetallic impurities.Among the fundamental studies that have determined the methodology of investigations of al model ,10 in wh217 and SmZ have been logically and widely used by investigators not only for the generic designation of alloys and magnets based on them, but also for plotted dependences “composition–physical properties”.In this context, the formulas Sm1-XZrXZ suggested by researchers of the Department of Magnetism, Tver State University supervised by Prof. Dmitry D. Mishin, Ph.D  is pnBm intermetallics is given in the form of n:mH or n:mR, where H and R indicate the hexagonal and rhombohedral crystal lattices, respectively).Currently, it is obvious that during heat treatment for the high-coercivity state, at the end of the solid solution heat treatment (SSHT), the starting matrix, in the classical sense of this term, is formed based on the disordered 1:7H phase ,4,5,6 (f7-type phase (space group P6/mmm), rather than 2:17R, which is typical of 1:5H and 2:17H [3/mmc [2Co17 phase. It is obtained by deformation of the SmCo5 structure through the disordered SmCo7 structure. In such a structure, Zr atoms occupy Sm sites.The 1:7H structure is the hexagonal TbCund 2:17H  or P63/mH [3/mmc . AccordiH [3/mmc , the hex2Ni17-type structure refined by the Rietveld method corresponds approximately to 0.5/0.5, 0.75/0.25, and 0.9/0.1 Sm/Zr or Co2/Zr population on sites 2b, 2d, and 4f/2c, respectively. Thus, the preferred sites for accepted Zr atoms are the Sm atom sites [The substitution of Zr atoms in the close hexagonal Thom sites .172 structure indicated that Zr atoms reside only on the 2b site [The analogous situation is characteristic for the Fe–Gd–Zr system. The Rietveld refinement of the Fe 2b site ,16.1-XZrXZ.Thus, the matrix being the initial solid solution, with good reason, is written in the form of Smn-1Co5n-1 homological series has not been in doubt for a long time. As for the 2:17R phase, Zr should definitely be included in parentheses together with the Co or 3d-metal mixture, Sm217. However, this phase, despite its high final volume in the alloy and high role in the formation of magnetic properties, is the secondary product of phase transformations.The substitutions Sm–Zr in crystal lattices of phases of the Sm1-XZrXZ alloys are:-the separation of the effect of main elements qualitatively differing in the configurations of electron shells  on the structure and properties of the studied alloys; and-1-XZrXZ.during the formation of the alloy structure in the course of heat treatment, the matrix  is the disordered 1:7H solid solution and almost all intermediate and final phases formed in the course of phase transformations, except 2:17R, adequately correspond to the formula SmThus, the arguments for the original formula of the Smz materials for permanent magnets in the literature, which were obtained by modern experimental equipment that is perfect from the point of view of the methodological approach [There are many results of experimental investigations of approach ,5,6,7,8.Z sintered magnet samples using results of the local analysis of some phases [Results obtained by methods including high-resolution transmission electron microscopy (HRTEM), nanoprobe energy dispersive x-ray spectroscopy (EDXS), nano-beam diffraction (NBD), etc. have allowed one to estimate the fine structure of e phases ,5,6,8. HZ alloys and magnets based on them:Z alloys, which was noted in numerous studies, can be observed with optical resolutions [The first obvious level of heterogeneity of the olutions ,19,20,21Z alloys and magnets based on them is the scale of the so-called cellular structure, which has been repeatedly studied, and consist of cells and cell boundaries that are penetrated by thin plates along the basal planes.The second level of heterogeneity typical of both Z alloys and magnets based on them.The third level of heterogeneity is the structure of boundaries between cells, which obviously undergo phase transformations at all stages of thermal aging of the The methodology of this study was built taking into account the evidence of several levels of heterogeneity of the Z alloys.This study was devoted to regularities mainly typical of the first, the most rough, level of the heterogeneity of materials based on the Z alloys were prepared by high-frequency induction melting in an argon atmosphere at an excess pressure of 5–10% using Al2O3 crucibles  and individual components such as Sm, Zr, Co, Fe, and Cu of >99.9, >99.97, >99.98, >99.7, and >99.97 wt.% purity, respectively. After melting, the alloys were not cast in a mold but slowly cooled in the crucible, which was carefully thermally isolated from a water-cooled inductor. As a result, an ingot characterized by coarse grains 4–6 mm in average size was obtained. With respect to the magnetic state, each grain had a single easy magnetization axis (EMA). These samples could not be classified among traditional single crystals since they are not single-phase. However, all phases of one grain were strictly collinear; in other words, the EMAs of each phase within a grain were parallel to each other. For simplicity, from here on, we call these samples pseudo-single crystals.The 1-XZrX(Co1-a-bCuaFeb)Z samples of the experimental series is given in The chemical composition of the SmThe chemical analysis of alloys was performed by optical emission spectroscopy with inductively coupled plasma .The heat treatment conditions used were similar to those described in the literature ,5,6,7,8.Pseudo-single crystal samples of each of the compositions in as-cast and heat-treated (under different conditions) states were ground to form balls 2.5–3.5 mm in diameter by grinding with gaseous nitrogen supplied under a pressure between two abrasive caps.MAX = 30 kOe). To measure the major hysteresis loop, the samples were preliminarily magnetized in a magnetic field  of no less than 100 kOe. The VSM was graduated using a Ni standard with σSNi = 54.4 G × cm3/g. A VSM-option of PPMS-9T Quantum Design installation was used to study the magnetic characteristics of samples, the coercive force of which at room temperature is above 30 kOe.The magnetic measurements of major and minor hysteresis loops of pseudo-single crystal spherical samples were performed at room temperature using a vibrating sample magnetometer   . The compositions of structural elements were studied by local x-ray energy dispersive analysis using pure elements as the standards. The EDX analysis of each of samples started from the determination of the integral chemical composition for the maximally possible section area in order to control the correspondence of the composition of the sample to the chemical analysis data for the associated series. Analogously, individual structural components were studied using areas knowingly apart from the interphase boundaries. Element scanning across the boundaries of the structural components was also used.Anisotropic sintered powder samples were prepared from these alloys using a standard powder metallurgy technique, which includes crashing in a nitrogen atmosphere, and milling in a vibrating ball mill. Compaction of powders in a transverse magnetic field, cold isostatic pressing, sintering in a hydrogen atmosphere, and heat treatment in an argon atmosphere were performed in accordance with regimes described above. The magnetic properties of the sintered samples (Ø15 × 7 mm) were measured in fields of 30 kOe using a completely closed magnetic circuit and an automatic recording flux meter .0.87Zr0.13(Co0.690Cu0.070Fe0.240)6.5 sample, which are typical of all samples under study.The given hysteresis loop is sufficiently wide for detailing the minor hysteresis loops and is also almost completely within the operating fields of VSM.The bottom curves correspond to the magnetization curves of sample demagnetized by reversing field and the series of minor hysteresis loops measured with progressively increased magnetized field. The top curves indicate magnetization curves of the same sample demagnetized with the alternating field with decreasing amplitude and the series of minor hysteresis loops measured with progressively increased magnetized field.This sample and samples of all alloys under study indicate the classic magnetic hysteresis related to the domain wall pinning mechanism.S = 116 G × cm3/g, which corresponds to values 4πJS = 12.5 kG and (BH)MAX = 39 MGOe.It can be seen that from the viewpoint of maximum energy product, the hysteresis loop is ultimate. The decrease in the magnetization, with allowance for the demagnetizing factor, starts in negative fields of ~8–9 kOe; the specific saturation magnetization of the sample was σThe hysteresis loops of the other samples are qualitatively similar to that given in 0.85Zr0.15(Co0.702Cu0.088Fe0.210)Z, and samples of series nos. 1–4 with the fixed z, Sm1-XZrX(Co0.702Cu0.088Fe0.210)6.4, are given in The magnetization reversal portions of the major hysteresis loops of samples of series no. 2  with the0.85Zr0.15(Co0.702Cu0.088Fe0.210)z alloys , the clearer the “shoulder”. Moreover, as the relative zirconium concentration (x) increases, the z value corresponding to the worsening loop squareness shifts to the low values. This can be clearly observed in comparing the dependences in 1-XZrXz alloys were qualitatively similar to those of the pseudo-single crystal samples. As an example, The dependences of the magnetic properties of the sintered powder samples on the chemical composition of the SmThe main difference of the sintered powder samples consists of the fact that the characteristic “shoulder” in the magnetization reversal curves of powder samples appeared for the compositions with z lower by 0.15–0.2 than that for the pseudo-single crystal samples of analogous compositions from all series of alloys under study. It is likely that the decrease in the threshold of the appearance of the “shoulder” is related to samarium and zirconium losses due to the oxidation and to the reactions with nonmetallic impurities and samarium losses due to evaporation during the course of manufacturing the sintered samples.0.85Zr0.15(Co0.702Cu0.088Fe0.210)6.4 alloy in the (a) as-cast state, (b) after solid solution heat treatment, and (c) after complete cycle of heat treatment for the high-coercivity state, which are typical of all series of samples.The microstructure of all samples under study consisted of three structural phase components (SPCs): dark-grey SPC ; bright-grey dendritic-like SPC (B), and almost white anisotropic SPC elongated along the basal plane of sample (C), which, for all samples, was localized within SPC A. The heat treatment did not lead to any substantive changes in the quantitative relations of structural components in all studied samples with the same chemical composition. After solid solution heat treatment and quenching, the characteristic composition contrast in the section disappeared; however, the cast prehistory of the cast state was clearly observed. After a complete heat treatment cycle, the contrast recovered but differed in the halftone pattern.A portion of the microstructure in The same colors were used in 6.3; in SPC B, 7.1; and in SPC C, 3.7 7.5, Sm8.4, and Sm12, respectively). Thus, the structural components were formed based on the 1:5 (A), 2:17 (B), and 2:7 (C) phases. This corresponds to results reported by Kianvash et al. [The energy-dispersive X-ray (EDX) microanalysis performed for the samples of all series of alloys in the high-coercivity state showed the similarity of all chemical compositions of the structural components. It is possible to state that the intermediate compositions corresponded to the common formulas: in SPC A, 1-XZrX(Co0.702Cu0.088Fe0.210)z alloys with the same relation of 3d elements (series 1–4).The analysis of the dependences given in As the relative Zr content (x) increases, the volume percentage of the structural component C increases. In this case, for each of the series of alloys, as the content of 3d elements (z) increased, the volume percentage of structural component C decreased.The volume percentages of structural components A and B corresponding to more than 90% of the total alloy volume changed monotonously  from the dominating volume percentage of structural component A to the dominating volume percentage of structural component B.A = VB), an intense increase in the coercive force was observed; after that (at VA < VB), the increase in HCJ decelerated and the dependence reached the horizontal “plateau”.As the relative content of 3d elements (z) and the volume percentage of structural component B increased in compositions corresponding to the equality of the volumes of structural components A and B (V1-XZrX(Co0.702Cu0.088Fe0.210)Z alloys, the transition through the chemical composition corresponding to the equality of volume contents of structural components A and B (VA = VB) was accompanied by the appearance of a “shoulder” in the magnetization reversal curves along with the end to increase in HCJ. Thus, the reaching of the alloy chemical compositions corresponding to the condition VB > VA was accompanied by the decrease in the squareness of hysteresis loops. This was satisfied for all series of Sm1-XZrXZ alloys.The comparison of the data given in 0.85Zr0.15(Co0.665Cu0.075Fe0.260)Z alloy. Thus, the following conclusion can be inferred.1-XZrXZ alloys in the high-coercivity state, which are characterized by the maximum  coercive force (HCJ) along with the high squareness of the loop, are reached in the case of equal volumes of structural components A and B comprising the main volume of the alloy and largely determining the hysteretic properties of the alloy.The major hysteresis loops for the SmThis conclusion also fully corresponds to samples of sintered magnets produced from the powders of the alloys in the experimental series Z alloys and is directly adjacent to the homogeneity range of the 2:17 phase in the 2:17H modification with the Th2Ni17-type structure Z samples under study in the as-cast and heat treated states shows that, for all of the studied compositions Z compositions, the structural formation at the side of the 2:7 phase is also of importance. To refine the processes that occurred, we considered the previously published results.However, in our opinion, to completely present the alloy formation of the R1-XZrXZ compositions.For the first time, the existence of the strict correspondence between the alloy composition, alloy microstructure , and hysteretic properties, which is given in 1-XZrXZ alloys can be easily etched and are more in contrast when compared to that of the Sm alloys because of the more simple set and configuration of phases observed for the actual composition range. In particular, the 5:19 phase was absent in the Gd–Co system. In the rare-earth series, when the atomic number increased, the atomic radius decreased and the 2:17H structure was stabilized, in particular, when Zr atoms were substituted for R atoms [It should be noted that the Gd-containing alloys were found to be preferred as model alloys for the investigation of composition–property correlations. The ferrimagnetic coupling of rare-earth and transition metal sublattices of the main phases decreased the magnetic moment of the samples and led to lower anisotropy field and, therefore, coercive force values. This is convenient for magnetic measurements in the limited range of quasi-static external magnetic fields. Moreover, the oxidation of Gd-containing alloys is lower than that of Sm-containing alloys, and the vapor pressure of Gd at 1000 °C is lower than that of Sm by seven orders of magnitude. These circumstances facilitate the preparation of samples. The microstructure of the Gd R atoms .It should be noted that Khan  had alre1-XZrXZ alloys Z alloys in the  corner of the quasi-ternary –Gd-Zr system Z alloys were similar to those for the Sm1-XZrXZ alloys Z alloys consists of the fact that because of the peculiarities noted for the Gd alloys, the dependence of the coercive force on the composition for all series exhibited a clear extremal behavior. The maximum of the dependence of coercive force strictly corresponds to the equality of the volumes of structural components A and B.The difference between the data given for the All the conclusions formulated by us based on the results obtained for the Sm-containing series of alloys were satisfied for the Gd-containing alloys.Z alloys.However, the extremal behavior of the dependences of the coercive force on the chemical and phase compositions for the Gd alloys allowed us to made an additional conclusion formulated for the Z alloys, a trend was clearly observed to broaden the coercivity peak (decrease in sharpness) with increasing Zr content and trend to shifting the peak to the low relative fraction of 3d-elements (z).For the experimental series of Z samples in the high-coercivity state tate see  and pseuis loops .0.85Zr0.15(Co0.70Cu0.09Fe0.21)z alloy series, the excess of the volume content of component B over component A (VB > VA) is accompanied by the decrease in the coercivity and, like in the case of Z alloys, by the progressive decrease in the squareness of the hysteresis loop.The comparison of CJ), localization of Zr in crystal lattices of Smn-1Co5n-1 phases, and the above constructed Co–Sm–Zr and –Sm–Zr phase diagrams.We calculated the phase relations for the –Gd–Zr system based on results given in The results are given in z alloys, obvious melting on the edge of the structural component C (2:7) was observed. Similar melting was observed by Kianvash et al. [This sketch of the diagram corresponds to temperatures slightly below the upper interval of SSHT temperature, at which, according to results of the analysis of microstructures of the h et al. .1-XZrXZ alloys Z alloys.Taking into account the data given in Z alloys  easily tCalculations show that the transition from the constructions performed for the Gd-containing compositions to the constructions for the Sm-containing compositions was mainly reached by elongation of the diagram in 1-XZrXZ alloys (used for manufacturing of permanent magnets) in the temperature range of SSHT, is given in The obtained sketch of the diagram, which indicates the phase relations in the SmZ system differed substantially from the diagrams given in [As can be seen from given in ,29,30,31given in ,23,24.As the Zr content in the alloys increases, the region of the high-temperature 1:7 phase, on the way out the “corridor” of the wide homogeneity range of the derivative 1:(5 + x) phase, “runs” in two directions corresponding to the back course of branches of the primary solidification of the alloy from liquid.P6/mmm). The second is the direction toward the homogeneity range of high-temperature hexagonal polymorphous modification of the 2:7 phase (space group 3/mmcP6).The first is the direction toward the homogeneity range of the high-temperature hexagonal polymorphous modification of the 2:17 phase (space group 5+X and Sm2Co17(h) are isostructural . It means that we should obtain a homogeneous region from SmCo5+X to Sm2Co17(h) in the phase diagram, at least, at temperatures near to the solidus. What we experimentally obtain is a eutectic in the neighborhood of SmCo7. An explanation for this discrepancy is still being sought” [The logicality, at least, of the first similar closure of high-temperature hexagonal structures was indicated repeatedly by Khan in the 18 studies related to the R–Co systems. In particular, he wrote: “It should be noted that SmCo sought” .Thus, at high temperatures, the 1:7-based solid solution range in the quasi-ternary phase diagram is present in the form of an asymmetric biconvex (toward the decreasing temperature) lens adjacent to the high-temperature homogeneity ranges of the 2:17 and 2:7 phases in their wider ranges.1-XZrX(Co0.70Cu0.09Fe0.21)Z.It should be noted that the given sketches of the phase diagrams are related to certain relations of 3d elements, which correspond to the formula RZ alloys in the concentration ranges promising for the production of high-coercivity permanent magnets. These findings were based on a quantitative analysis of the microstructures of the first, most crude level of heterogeneity of these alloys in a highly coercive state.We present here an original view on the issues of the structural formation of the 1-XZrXZ is physically most reasonable from the viewpoint of the effect of main elements with the qualitatively different electron shells  on the structure and properties of the studied alloys, and with allowance for the fact that the initial structural state is the matrix being the single-phase disordered 1:7H solid solution, and almost all intermediate and final phases (except the 2:17R phase) formed in the course of phase transformations corresponded better to the formula R1-XZrXZ.The formula of the alloys R1-XZrXZ alloys in optical resolutions was formed by three structural components based on the 1:5 (A), 2:17 (B), and 2:7 (C) phases; the qualitative relationships A:B:C varied within wide ranges as the chemical composition changed monotonously.The microstructure of the R1-XZrXZ alloys of all experimental series, the increase in the relative content of 3d elements (z) was accompanied by monotonic changing of the volumes of structural components A and B, corresponding to 90% of the total alloy volume from the dominant volume of structural component A to the dominant volume of structural component B.For the studied composition ranges of the Sm1-XZrXZ alloys in the high-coercivity state were strictly controlled by the volume percentage ratio of the structural components A and B.The hysteretic properties of the SmA = VB corresponded to the combination of the high coercive force (HCJ) and ultimate squareness of the hysteresis loop; the predominance of the volume of component B was accompanied by the decrease in the squareness of the hysteresis loop.The dominant volume of the structural component A in the alloy ensured the ultimate squareness of the hysteresis loop of the samples; the equality of the volumes V1-XZrXZ alloys in the high-coercivity state was accompanied by the decrease in the precision of the dependences of the coercive force on z and the shift of alloy compositions with the equal volumes VA = VB toward the lower relative content of 3d-elements (z).The increase in the relative Zr content (x) in the SmAt temperatures corresponding to the supersaturated solid solution treatment, the section of the ternary diagram of the Co–Sm–Zr system undergoes a significant transformation upon introduction of Cu and Fe into the alloy and the transition to the –Sm–Zr quasi-ternary section.In the quasi-ternary phase diagram of –Sm–Zr, at SSHT temperatures, the region of the 1:7-phase-based solid solution had the form of an asymmetric lens, biconvex toward lower temperatures; the edges of the lens were adjacent to the homogeneity ranges of the 2:17 and 2:7 phases in their widest, high-temperature regions.It was shown that:Z alloys in high-coercivity state for composition ranges that are of importance in the manufacture of permanent magnets.In future studies, we will consider the peculiarities of magnetization reversal and the composition of the structural components of the


The subject of the research is a polymer composite with a matrix base of epoxy resin L285 cured with H285 hardener, and a physical modifier of friction in the form of alundum. The article presents an analysis of findings of tribological examinations. The authors evaluated the influence of the modifier properties in the form of alundum, i.e., mass share and grain size, on the abrasive wear of a composite, defined as loss of weight as well as on roughness parameters and selected mechanical properties. The tribological examinations have been extended by measurements of hardness and density of the prepared composites. The obtained results of tribological examinations showed an increase in the average value of weight loss in relation to the loss of sample weight loss between the cycles. The influence of both the grain size and the mass percentage share of alundum upon the increase in the longitudinal modulus of elasticity was also observed. On the basis of the obtained results, it was found that alundum of grain sizes equal to F220 and F240 exerted the best influence on the reduction of abrasive wear of the tested samples. In the case of F220, it was 14.04% of the average value of the weight loss between the cycles for all percentage shares of the used grains. Modern industries, especially mechanical engineering  and the –good constructional properties,–small specific gravity,–easy moulding of products ,–diversity of processing techniques,–possibility of differentiating the properties by modifying the use of intermediate inputs and processing techniques.Composite materials have long been used due to ,14:–gooIn addition, their use was popularized by the fact that it was possible to introduce changes to the structural properties by including modifiers in their composition ,16.–natural organic fillers ,–inorganic fillers,–synthetic fillers .When classifying modifiers by origin, it is possible to distinguish ,18,19,20–powder fillers ,–fibre fillers .Taking into consideration the form in which the fillers appear, the following can be distinguished:Physical modifiers and reinforcing materials improve material properties such as density, hardness, impact strength, and structural properties, including the tribological ones ,22.–wear rate,–linear or volumetric wear,–force of friction,–coefficient of friction,–temperature of the contact area.The problems of tribological wear and tear of machine and equipment components ultimately translate into their performance and operational reliability. The occurrence of numerous machine operational units makes it necessary to limit this irreversible phenomenon, which is achieved, for example, by using construction materials with improved anti-wear properties. Conducting tribological research in industrial conditions is difficult and costly . TherefoA literature analysis related to tribological issues repeatedly refers to automotive-related research ,25. The Aeronautics is one of the areas of application of construction materials in which tribological aspects are subject to an in-depth analysis. This is influenced by specific operating conditions of manufactured components combined with high safety requirements. The available literature presents the research results on selected areas that require improvement in friction parameters. They relate, among others, to aircraft brakes, e.g., of the carbon-carbon type (research into the effect of infiltration of silicon carbide nano-additives on a porous composite used to manufacture brake discs ), contaiIn addition to the areas indicated, the authors of the research publish the obtained findings with regard to improved properties of aircraft skin. In this case, it is important to take into account a wide range of forces occurring in operation, also including extreme operating conditions which result from humidity changes depending on a changing altitude, a wide range of temperature values, and the occurrence of abrasive wear caused by electrostatic attraction of dust. Improvement in tribological properties of composite materials can be conducted by plasma spraying  or the uThe research part of this publication deals with an inorganic powder filler in the form of an abrasive, i.e., alundum of different grain sizes. The granularity of abrasives is determined on the basis of the sieve analysis (up to 53 μm). The digital marking used in the notation indicates the number of holes found in a specified unit of the sieve area. The thickest grains are identified by the symbol F8, whereas the finest ones by F1200. Apart from the characteristics of the abrasives in the form of described grain, their hardness is often taken into consideration. For example, granite has the hardness of approximately 14 GPa, quartz of approximately 11 GPa, while the hardness of alundum equals 20–22 GPa on the Vickers scale .In the available literature, it is possible to find numerous examples of research conducted so far in which an influence of various fillers that acted as friction modifiers was determined. The authors checked the influence of various mass percentage shares of graphite on tribological properties ,38,39,402 [2O3 [In the available literature, among other friction modifiers, it is possible to find, e.g.: polyphenylene sulphide (PPS) , TiO2 [42 ,46, or A2 [2O3 . Researc2 [2O3  such as 2 [2O3 , sugarca2 [2O3 , or jute2 [2O3 . It is w2 [2O3 .The cited publications, which present only a narrow slice of the variety of areas of tribological research, prove that the search for dependencies describing the processes of material consumption is still up to date. The observed intensification of the development of materials engineering, apart from defining the obtained mechanical properties for a given material, also enforces determining tribological properties, taking into account further possibilities of exploiting these materials in construction solutions, which include friction pairs. It is, therefore, appropriate to identify the existing dependencies that may, in the long term, determine the application of new solutions in practice or indicate operational recommendations.2O3, 99% pure, and with F220, F240, F280, F320, F360 grains (the division of abrasive grain in accordance with FEPA 42-2:2006). The actual grain size, in accordance with the standard FEPA is, respectively: 53, 44.5, 36.5, 29.2, and 22.8 µm [The research was carried out using polymer composites with a matrix base, whose reinforcement phase was alundum (EA symbol) AlThe designation of the fraction as F0 indicates pure epoxy resin, without the addition of alundum. The examinations were made in the laboratory of the Department of Airframe and Engine at the Military University of Aviation in Dęblin (Poland). For the preparation of samples by gravity casting, epoxy resin with the trade name L285, along with H285 MGS hardener, were used, as well as alundum with mass percentage share of 5%, 10%, 15%, 20%, and 25%, respectively, for each of the mentioned grains. Gravity casting consisted of preparing a mixture of resin, hardener, and EA, and then filling the moulds with it. After removing the air bubbles using ultrasonic waves, the mould was allowed to stand for 24 h at room temperature until the resin was completely cured. Each batch contained 10 samples and the results were averaged. In order to determine whether the grain distribution was homogeneous, the fracture surface was observed on the optical microscope OLIMPUS BX53M+ . The produced samples were tested in conditions of abrasive wear in a reciprocating motion using the tribotester—Taber Linear Abraser, model 5750 . The abrasive stone, used on the test stand, was 6.6 mm in diameter. Its grain equalled 200. The friction path was 177.8 mm/cycle and the total load equalled 1850 g. Prior to the friction process, the samples were weighed. Additionally, their initial weight was determined. After 100, 300, 600, and 1000 cycles, the weight of the samples was determined again using the precision laboratory balance XSE205DU/M . After each measurement, the sample was carefully cleaned, also removing the products of friction. Moreover, in accordance with the manufacturer’s recommendations, the abrasive stone and the sample were cleaned. After each interval of cycles, the authors measured the selected roughness parameters in a cross-section, perpendicular to the direction of the path covered by an anti-sample. The measurement of the roughness parameters, Ra  and Rmax (roughness depth), was made by means of an optical profilometer . Moreover, in accordance with the PN-EN ISO 527-2:2012  norm, a In the further part of the research programme, the samples were subjected to hardness measurements in accordance with the Shore D method at a three-second load of 5000 g, in compliance with the PN-ISO 868  norm usi2 > 0.8 can be observed, which indicates good correspondence of the regression function to the obtained test results. Moreover, all the results, regardless of the type of the grain size and the amount of the filler, due to the similarity of the nature of the discrepancy, indicate high repeatability and consequently the predictability of weight losses during the friction process.Moreover, the conducted observations focused on determining whether there is an influence of the grain size of alundum on the weight loss at a given percentage. Thus, for each percentage mass share of the physical friction modifier, a graph was generated showing weight loss after 0, 100, 300, 600, and 1000 cycles .The generated graphs of weight loss from the initial state for particular percentage shares of physical friction modifiers proved that in the majority of cases , the greatest weight loss was observed in samples with the presence of alundum whose grain equaled F360. It was similar to the weight loss of pure resin samples. This results from the fact that the sharp-edged alundum grains . The biggest difference against the reference samples was observed for the EA F360 composite, amounting to 2788.44 MPa . Thereby, the smallest difference with regard to the reference samples was observed in the case of 5% mass share of EA F220 in the composite.On the basis of On the basis of the conducted research, it was found that along with a rise of the average value of weight loss, the weight loss between the cycles was also growing. When comparing composite wear with an addition of various grains, it was found that alundum of grain sizes equal to F220 and F240 exerted the biggest influence on the reduction of abrasive wear of the tested samples. However, taking into account the roughness parameters, the F280 alundum had the effect of obtaining the lowest roughness parameters, Ra approximately at level 1 µm, and Rmax averaged 76 µm for a mass share of 15%. Regardless of the mass share for these grains, the abrasive wear was the lowest after each cycle interval . The addition of a physical modifier of friction influenced the selected mechanical properties of the composites, both in the case of the value of the modulus of longitudinal elasticity and the max principal stresses value. The highest value of the material’s modulus of longitudinal elasticity was obtained for 25% of the mass share of alundum. Taking into account the differences against the reference samples, it was observed that for the EA F360 composite, it is the largest, amounting to 2788.44 MPa, with an increase in the material’s modulus of longitudinal elasticity by 26% in relation to samples made of pure resin. The smallest effect of the presence of the additive against the reference samples was observed in the case of 10% mass share of EA F220 in the composite (an increase by merely 0.7%). However, no relationships were observed between them and the obtained hardness of the composites. The obtained results indicate a lack of linear relationships occurring between the observed properties and the composition of the composites. Currently, the authors are conducting research aimed at making multi-criteria optimisation of composites’ composition containing alundum with epoxy matrix base.


Resistance training (RT) under hypoxic conditions has been used to increase muscular performance under normoxic conditions in young people. However, the effects of RT and thus of RT under hypoxia (RTH) could also be valuable for parameters of physical capacity and body composition across the lifespan. Therefore, we compared the effects of low- to moderate-load RTH with matched designed RT on muscular strength capacity, cardiopulmonary capacity, hematological adaptation, and body composition in young and older people.2): ∼80–85%. Changes and differences in maximal isokinetic strength, cardiopulmonary capacity, total hemoglobin mass (tHb), blood volume (BV), fat free mass (FFM), and fat mass (FM) were determined pre–post, and the acute reaction of erythropoietin (EPO) was tested during the intervention.In a pre–post randomized, blinded, and controlled experiment, 42 young (18 to 30 year) and 42 older (60 to 75 year) participants were randomly assigned to RTH or RT . Both groups performed eight resistance exercises  four times a week over 5 weeks. The intensity of hypoxic air for the RTH was administered individually in regards to the oxygen saturation of the blood (SpOp = 0.120 to 1.000) were found between RTH and RT within the age groups. However, within the four groups, isolated significant improvements (p < 0.050) of the single groups were observed regarding the muscular strength of the legs and the cardiopulmonary capacity.In all parameters, no significant pre–post differences in mean changes . Resistance training is usually used to increase or maintain muscular strength and muscle mass among all age groups. Since age-related loss of muscular strength and muscle mass are associated with an increased manifestation of a lower functional capability , which is why there is no evidence regarding its effects in older people. However, RTH could also be a valuable exercise intervention strategy with a higher effectiveness than a resistance training under normoxic conditions for older people. This assumption is based on study results and the associated discussion that hypoxia and hypoxia in combination with physical exercise  is a positive evaluated, auspicious intervention strategy against age-related changes of the physical and mental health , b as weTherefore, the aim of this study was to investigate the effect of RTH in comparison to matched designed resistance training under normoxia on muscular strength capacity, cardiopulmonary capacity, hematological parameters , and body composition [fat mass (FM) and fat free mass (FFM)] in young and older people.Regarding this aim, the design of RTH is decisive. With respect to the current recommendations for RTH, low to moderate loads [∼20 to 50% of the one-repetition maximum (1RM)], medium to high volume , and short inter-set pauses (∼30 s) are frequently recommended in order to reach a high level of metabolic stress and therefore to achieve a high stimuli for adaptations of the muscles , 2017. IA total of 42 young people (18 to 35 years) and 42 older people (60 to 75 years) were recruited using print advertisement in a regional newspaper and with flyers in regional sport and health facilities. To participate in this study, the participants had to fulfill the following requirements: non-smoker; not pregnant; no disorders of the musculoskeletal system, the cardio-pulmonary system, the kidneys, and the brain; no blood donation or loss of blood (>200 ml) in the last 3 months; no sojourn in altitude (>1800 m) in the last 3 months; no additional physical activities other than those during the period of the study. Moreover, the participants were instructed to maintain their regular dietary consumption, to avoid intense exercises, and not to start any nutritional supplementation during the entire study period.®, Germany) assigned to one of the following two groups: resistance training under normobaric hypoxia  or resistance training under normoxic conditions . The intervention period was followed by a 2-week post-diagnostic period , twice each with young and older people to include 42 young and 42 older subjects in this study. This was necessary because the planned intervention was limited to the amount of altitude generators and strength training equipment available. Each block started with a recruitment period of 4 weeks. At the end of this period, the participants were invited to visit an information session where the aim and experimental procedure were explained and written informed consent was obtained from each subject. Subsequently, a 2-week pre-diagnostic period and then a 5-week intervention period were conducted. Prior to the start of the pre-diagnostic period, the young and older participants were randomly  and the staff conducting the diagnostics and analyzing the data was also blinded (double blinded study). The study was officially approved and authorized by the Ethical Committee of the Otto von Guericke University Magdeburg, Germany (Nr. 74/15).®, United States), the participants did not know their group affiliation (blinded conditions). Each intervention session had a total duration of 3 h and consisted of three phases lasting 60 min each.During the 5-week intervention period, the RTH and RT were conducted four times per week  within one of three encounter groups each comprised of seven participants . Prior the start of the study, the participants were able to choose one of these encounter groups in which they were asked to train continuously afterward. Thus, the participants of the RTH and RT were mixed in the three encounter groups. Since the hypoxic air (for the RTH) as well as the normoxic air (for the RT) were individually applied via a facemask  through an altitude generator  was individually adjusted to the oxygen saturation of the blood (SpO2) for the RTH groups. For the acclimatization to the hypoxic conditions in the RTH groups, the SpO2 was gradually reduced: first week 85 to 88%, second to fifth week 80 to 83%. The SpO2 was checked every 5–10 min with a pulse oximeter  and the FiO2 was adjusted if the SpO2 did not meet the predefined SpO2-values . The RT groups received air with a FiO2 of 20.9%. To ensure the blinded conditions, the SpO2 was measured in the RT groups like in the RTH groups and sham adjustments were undertaken with the altitude generators. Phase 2 was comprised of supervised resistance training (RTH/RT) with whole-body resistance exercises on resistance training equipment by Proxomed . Before the resistance training started, a supervised warm-up program of 5 min (dynamic stretching for the major muscle groups) was performed by the participants. To acclimatize to the hypoxic condition, the FiO2 was gradually reduced and individually adjusted to a SpO2 between 85 and 90% in the first week and to a SpO2 between 80 and 85% from the second to the fifth week for the RTH groups. The RT groups continued to receive air with a FiO2 of 20.9%. The SpO2 was checked in both groups after each exercise and adjusted like in phases 1 and 3.During phases 1 and 3, the participants breathed hypoxic air [phase of intermittent hypoxic exposure (IHE) for the RTH groups] or normoxic air (rest phase for the RT groups) according to their group affiliation in a seated position. With regard to the recommendation by The resistance training program consisted of two training plans (A and B) that were performed alternately. Training plans A and B each included multiple-set circuit training with eight machine-based resistance exercises: Training plan A: (i) lat pulldown, (ii) butterfly reverse, (iii) upper body extension, (iv) knee flexion, (v) vertical row, (vi) upper body rotation, (vii) biceps curl left, (viii) biceps curl right; Training plan B: (i) triceps press, (ii) leg press, (iii) butter fly, (iv) upper body flexion, (v) knee extension, (vi) chest press, (vii) shoulder side rises left, and (viii) shoulder side rises right. The order of performing the exercises was the same for each participant, but due to the fact that all participants trained at the same time, each one had a different start-exercise and therefore an individual sequence. Each exercise of the training plans was performed with three sets of 15 repetitions, a cadence of 2-0-2-1 s (a metronome with 60 beats per minute was used), an inter-set pause of 30 s, with full range of motion, and with a load that was individually adjusted to a subjective perceived exertion using the CR-10 scale  on a bicycle ergometer. An isokinetic maximal strength test and a body composition test as well as the determination of total hemoglobin mass (tHb) and blood volume (BV) were performed at the second visit. The participants completed a questionnaire on physical activity and got an introduction to the training plan at their third visit see .–1) on a bicycle ergometer  in combination with spirometry . The test was performed until subjective maximal physical exhaustion. During the whole test, vital parameters  and heart rhythm  were observed by a medical doctor. Blood pressure was measured , oxygen consumption at 1 watt/kg body mass (VO21W/kg), maximal power output (POmax), power output at the second lactate threshold (POLT2) , and the systolic and diastolic blood pressure  at rest and at these watt-stages which were performed by all participants.The cardiopulmonary capacity was tested with an incremental load test  was tested for the exercises elbow flexion and elbow extension as well as knee extension and knee flexion for the dominant and non-dominant extremities using the BTE PrimusRS . The test procedure and the data processing were conducted following the recommendations by max was first tested for the lower extremities and then for the upper extremities. The test procedure started with a 5-min warm-up for the lower extremities on a bicycle ergometer  and for the upper extremities on an arm crank ergometer  with a load of 1 W/kg and 0.5 W/kg (each with a cadence of ∼70 min–1). The test procedure was carried out as follows: warm-up for the lower extremities, Fmax test dominant side for the exercise knee extension and flexion, 3 min rest, Fmax test non-dominant side for the exercise knee extension and flexion, warm-up for the upper extremities, Fmax test dominant side for the exercise elbow extension and flexion, 3 min rest, Fmax test non-dominant side for the exercise elbow extension and flexion. The Fmax tests were performed with two familiarization sets of five repetitions . Thereafter, the Fmax tests were performed with five repetitions. The movement velocity was 60°/s in each set , fat free mass (FFM)] was determined with the bioelectrical impedance analysis test BIA;  using Nust, 9th, 11th, 13th, and 15th minute during the rebreathing period to analyze hemoglobin, hematocrit, and the fraction of carboxyhemoglobin .The hematological parameters tHb and BV were determined by the carbon monoxide (CO) rebreathing method by basic), extracurricular activity (FFkAextracurr), sports activity (FFkAsports), and the sum of the total activity . To monitor the physical activity during the intervention period, the participants received an accelerometer  which was worn on the left wrist. Due to the maximum recording period of the GT3X of 3 weeks, just the first 3 weeks of the intervention period were monitored. The GT3X calculated the kilocalorie per day  and the metabolic equivalent per day (MET/d). Wearing of the accelerometer was elective for the participants and a total of 10 participants decided not to wear the GT3X . Because it is known that physical performance and adaptation processes are related to physical activity . The FFkA determines the health-related physical activity retrospective for 1 week (in hours per week) . Here, aactivity , we used®. The pre-diagnostic methods of visit one and two were adequately conducted in the post-diagnostic (see The introduction phase in the strength training plan aimed to familiarize the participants with the training procedure and included: (i) the individual determination of an adequate anthropometric setting of the strength training equipment, (ii) the explanation of the execution of each exercise, and (iii) the initial determination of the load for each resistance exercise for a subjectively perceived effort of 7 on the CR-10 scale. These were conducted by a sport scientist who was licensed in exercise training for public health, prevention, and rehabilitation and attended a course for using the strength training equipment by Proxomedstic see .nd , 11th , and 19th  of the 20 training sessions to determine the acute reaction  of the hematopoietic hormone erythropoietin (EPO) (determined in the blood serum). Moreover, ferritin and the glomerular filtration rate (GFR) were determined in t1 to check the iron level and renal function  and arms  during the 20ning see . Here weThe data was checked for normal distribution using the Kolmogorov-Smirnov-Test. For normally distributed data, the mean and standard deviation and for non-normally distributed data the median and the interquartile range is reported.t-test or Mann–Whitney U-test (group effect). To check the homogeneous distribution of the sex within and across groups, we conducted a chi-squared test. In a primary analysis, the effect of the RTH vs. the RT on the parameters of the cardiopulmonary capacity, the muscular strength capacity, EPO, the hematological parameter tHb and BV, and the body composition  were analyzed by performing a two-way ANOVA with repeated measures (time × group effect). Since multiple tests can lead to false-positive results .To compare the loads as percentage of the 1RM between the four groups, a one-way ANOVA without repeated measures (main group effect) was performed. When a significant effect was identified, the Bonferroni p < 0.05 for all analyses. The data processing and analyses were done using the software Excel  and SPSS Statistics .In conjunction with the time × group effects, the effect size partial eta squared  were similar between the groups RTH and RT within the age cohorts . The samp = 0.010).Gender distribution was homogeneous within the age-groups , there were no significant differences for the used loads within the resistance exercises between the four groups. The difference of the used loads between the four groups was maximal 8% (RTH young vs. RT old for the exercise knee extension) and usually only between 2 and 5%.The values of the used loads in the groups RTH and RT of the young and older participants are shown in 2) was comparable over the 5-week training intervention period during the IHE phase but significantly lower during the RTH phase for the young age cohort in comparison to the older age cohort of the RTH groups  was also lower in the group RTH young than in the group RTH old  as well as at the 4th (p = 0.020) and 5th (p = 0.005) week during the RTH phase. From week 2 to 5  and during the RTH phase from week two to four , the SpO2 was significantly lower for the group RTH young in comparison to the group RTH old  in the group RT old only narrowly missed a significant effect  and 50 W (−10 mmHg: p = 0.012), 75 W (−10 mmHg: p = 0.001), and 100 W (−10 mmHg: p = 0.008) during the cardiopulmonary exercise test. The DBP was significantly higher from pre to post in the group RT old at 75 W (+ 10 mmHg: p = 0.030) and 100 W (+ 7.5 mmHg: p = 0.039)  see .We found no significant differences regarding the changes of the maximal isokinetic muscular strength between the groups RTH and RT within the young or older age cohorts  was not affected by the 5-week RTH or RT for the young or older subjects see , 3.Regarding the primary analysis, we found no superior effects of RTH in comparison to matched designed RT on physical performance (cardiopulmonary and muscle strength capacity), tHb, BV, and body composition . Heterogeneous results regarding the effect of this type of RTH are currently available: higher effectiveness of the RTH regarding the improvement of muscle hypertrophy , muscle Regarding the muscle strength capacity , we found only partially significant improvements (only for the muscles of the legs but not for the muscles of the arms with no differences between the groups RTH and RT within the age cohorts). The increase in the maximal isokinetic muscle strength of the legs was lower for both groups of the young people than in both groups of the older people. Considering the measurement error of the used device and method for maximal isokinetic strength testing  and POmax (+ 5.2%) as well as the significant decrease in the submaximal HR  in the group RTH of the older cohort suggest that the RTH may have had a positive effect on the cardiopulmonary system. However, there was also an improvement of the VO2peak of + 10.5% in the group RT of the older participants, which can be considered as practically relevant, although it was not statistically significant. In the young people, we observed a significantly improved exercise economy in the group RTH (VO21W/kg: −4.7%) and in the group RT (HR1W/kg: −6.6%). Therefore, it appears that the resistance training itself also had a positive effect on the cardiopulmonary system of the young and older people.Even though the cardiopulmonary capacity of the group RTH was not superiorly improved compared to the group RT in both age cohorts (no significant time × group effects), the significant increase in the VO2peak and/or the POmax between 1 and 8% in young people after several weeks of training (8 weeks to 6 months)  on the cardiopulmonary capacity (endurance performance) is positively associated with an increase of the VO2 months) . Compara months) . Therefopeak in young people. In contrast, the study by Additionally, previous studies that have investigated a low-load RTH , but without a significant effect. No changes or slight decreases of maximal −5 mmHg were reached in the groups RT young and RT old. These changes of the SBP correspond to the extent that can be expected by resistance training . The reduction in tHb in the RTH young (−2.5%) is not understandable for us because usually no change or a positive effect is associated with hypoxia  or exercBody composition, and in particular the FFM or lean body mass, was analyzed in previous RTH studies to indirectly research the effect of this type of intervention on muscle mass/muscular hypertrophy . We did 2) during the IHE and RTH phase met the predefined hypoxic intensity . Even though the differences in the SpO2 were only small (SpO2 2 to 3%), the SpO2 was still significantly lower in the young than in the older age cohort during the IHE and the RTH phase in the second to the fifth week. The hypoxic intensity is crucial regarding the strength performance  and might therefore also be crucial for adaptational processes. While 2 13 vs. 16%). However, both studies used different exercise variables. Thus, it can only be assumed that the hypoxic intensity can be decisive regarding strength performance, but does not have to be (depending on the design of the resistance training). In addition, 2 16%, SpO2 ∼86%) vs. high-level hypoxia  . Here, the difference between the internal hypoxic intensity was much higher than in our study. Although we cannot be certain, based on the small  differences in the internal hypoxic intensity in the current study, we are convinced that these differences were not practically crucial, because the stimuli for adaptation processes were different between young and older people.The internal hypoxic intensity . Given that gender can influence adaptational processes  and thusRTH used with low to moderate loads had no additional effects on the muscular strength capacity, cardiopulmonary capacity, EPO, tHb, blood volume, fat free mass, or fat mass in young people and older people. However, this statement can only be applied for RTH designed as in this study and after an intervention period of 5 weeks. Based on previous RTH-studies as well as on the knowledge from resistance training studies in general, it can be assumed that the expected higher effects can still be achieved with the change of exercise variables of RTH . Therefore, further RTH intervention studies (particularly those with longer intervention periods) are strongly recommended to reinforce knowledge about the adaptational processes and the effects of this type of resistance training in young and older people.The datasets generated for this study are available on request to the corresponding author.The studies involving human participants were reviewed and approved by Ethical Committee of the Otto von Guericke University Magdeburg/Germany (Nr. 74/15). The patients/participants provided their written informed consent to participate in this study.AT and LS contributed to the study design. AT contributed to the recruitment of participants, implementation of the intervention, data analysis, and manuscript drafting. AT and BP contributed to the diagnostic. LS and BP contributed to the review of the drafted manuscript.The authors declare that the research was conducted in the absence of any commercial or financial relationships that could be construed as a potential conflict of interest.


A concrete-filled steel tube (CFT) combines the advantages of concrete and steel in construction and structural applications. However, research on the time-dependent deformation of the CFT under long-term sustained loading are still limited, particularly for stress transfer between the steel tube and concrete due to creep. This study investigated the creep behavior of CFT over a long period of 400 days. The creep and shrinkage strain of CFT was significantly lower than those of concrete that was not confined within a steel tube. The vertical strains of the steel tube and concrete core were almost identical, and it was shown that they were well bonded and acted as a composite. The vertical stress of steel increased by 32.7%, whereas the vertical stress of concrete decreased by 15.8% at 375 days. The stress transfer is notable and cannot be neglected in CFT design. Moreover, the results of creep and shrinkage were compared to prediction values of the B4 model and B4-TW model to verify their validity. Concrete-filled steel tubes are composite structural members made of a hollow steel tube filled with concrete. These composites have been widely used in practical fields because of their high ductility, high strength, good fireproof properties, excellent seismic resistance, low creep and shrinkage benefits in comparison with reinforced concrete or steel structures ,2,3. NevMost experiments have analyzed the time-dependent behavior of the CFT with various affecting parameters, including structure shapes, types of concrete and steel tubes, and applied stress. The creep and shrinkage of square and rectangle section specimens were tested by Uy et al.  and Han Additionally, some researchers investigated the strain and stress of the steel tube and concrete core of the CFT member. Ichinose et al.  and Y. WThe CFT is made from two materials so that the composite has different properties and the time-dependent deformation is also more complex than separated components. Although some researchers have performed analyses of the properties of long-term behaviors, the time-dependent deformation of CFT under long-term loading has not been fully studied. Notably, a few studies investigated the stress transfer between the steel tube and concrete due to creep.Long term creep tests on the time-dependent deformation behavior of CFT using high strength self-compacting concrete (SCC) embedded with a vibrating wire strain gauge (VWSG) were conducted for 400 days. The observation and estimation of the stress transfer between the concrete core and steel tube under a sustained load were also provided. The stress transfer magnitude cannot be neglected and must be taken into consideration in CFT design. Moreover, the measured creep and shrinkage values were compared to prediction values of the B4 model and B4-TW model to verify their validity.2/kg. Mineral admixtures included ground-granulated blast-furnace slag (GGBS) and silica fume (SF) with a specific surface area of 435 m2/kg and 15,000 m2/kg, respectively. The used aggregates were fine, with fineness modulus of 3.1 and sandstone was used as a coarse aggregate (CA) with a nominal maximum size of 12 mm. To enhance the workability, a high-range water-reducing agent (type F) was used as the superplasticizer (SP) admixture. An expansive additive (EA) was incorporated into all specimens to reduce the shrinkage. The mix proportion of the concrete is presented in 3 , 20kg/cm3 , respectively.The cementitious materials used in the study were cement, slag and silica fume. Type II Portland cement complying with ASTM C150  was usedA total of six CFTs used cylindrical steel tubes with an in outer diameter of 108 mm, height of 300 mm and steel tube thickness of 3 mm. The steel tube had a yield strength of 250 MPa, grade satisfied ASTM A36. The CFT specimens had a diameter to thickness ratio (D/t) of 36, and a steel ratio (proportion of steel area to concrete area) of 12.1%. The composite members were not only carried out on basic creep and autogenous shrinkage tests but also used to determine the strain and stress behavior of the steel tube and concrete core. The CFT specimens were only used for basic creep and autogenous shrinkage tests because the role of the steel tube is to act as a permanent framework to maintain an enclosed environment so that concrete core in the steel tube is considered as a sealed specimen. The creep tests of the SCC specimens were conducted on Φ100 mm × 300 mm cylinders, while the shrinkage and length change tests of SCC specimens used two different sizes: Φ100 mm × 300 mm cylinders and 100 mm × 100 mm × 285 mm prisms. Each test set included three specimens. The specimens used in creep, shrinkage and length change tests are presented in The compressive strength and elastic modulus tests were conducted with cylindrical specimens of 100 mm by 200 mm dimensions. A total of 90 specimens were prepared to determine the mechanical properties at different times. Sixty specimens were used to investigate the strength and modulus of elasticity for ages of concrete at 7, 14, 28, 56 and 91 days. Another thirty cylindrical specimens were only evaluated for compressive strength in five creep tests at loading ages of 14 days and 28 days.All test specimens were prepared in the laboratory and cured in controlled conditions. The slump test of the SCC mixture was conducted by Abrams cone, according to ASTM C143 , while tThe compressive strength and the modulus of elasticity were tested according to ASTM C39  and ASTMζ is temperature-compensated strain; −6 mm/mm), respectively; K is coefficient of strain gauge, which is 0.98; The creep tests of the CFT and SCC specimens were conducted according to ASTM C512 . All timThe bottom end of each steel tube of the CFT specimens was welded to a steel base plate of 13 cm × 13 cm in cross section. After the VWSG was installed, the concrete was filled into the mold, and immediately after casting, a plastic sheet was coated on the exposed surface to prevent evaporation. At loading ages of 14 days and 28 days from the casting, the specimens were grinded smooth, as shown in While the CFT specimens were only subjected to a basic creep test, for the SCC specimens both basic creep and total creep were measured. The VWSG was also preinstalled in the mold of the SCC specimens and then the concrete was poured in. After 24 h from the casting, all the specimens were demolded and placed in moist curing conditions until the loading ages of 14 days and 28 days. After curing, the specimens of the total creep experiment were moved to be stored in a drying room with an automatically controlled temperature of 23 ± 2 °C and relative humidity of 50 ± 2%, whereas the basic creep specimens were s = kept in the moist environment  until the end of test. The SCC specimens were connected by high strength gypsum before being put into the loading frame as presented in The concrete core of the CFT specimens only experiences autogenous shrinkage as a result of the self-desiccation process, but the SCC specimens were assessed both in terms of total shrinkage and length change up to 400 days. All shrinkage specimens were demolded 24 h after casting and cured in a standardized chamber . When the concrete ages were cured for 14 days and 28 days, the C-SCC-TS14 and P-SCC-TS28 specimens—for which total shrinkage was determined—were stored in a drying room , respectively. Moreover, for the purpose of evaluation, the length change effects of the concrete on the creep data, the C-SCC-LW14 and P-SCC-LW28 specimens were placed in the same environment as the basic creep specimens at a designated time. The shrinkage strain of the cylindrical specimens can be recorded at an initial setting time (the beginning of solidification) by the VWSG. After the initial setting time of fresh concrete was determined based on ASTM C403 , the shrMany present models predict the creep and shrinkage of concrete, for example the ACI 209R-92 ,25, fib According to the B4-TW model, the creep compliance function is given byn is the empirical parameter, .In which 3), respectively; 3); factors With 3); The B4 model has different equation details in comparison with the B4-TW model. The creep compliance function is determined bycm) is replaced by the cement content (c).The According to the B4-TW model, the total shrinkage strain is given asV/S is the volume to surface ratio.In which The autogenous shrinkage of concrete without mineral admixture is expressed asAutogenous shrinkage of concrete with mineral admixtures is calculated byThe B4 model includes total shrinkage, expressed byOther shrinkage formulae of the B4 model are the same as in the case of the concrete without the mineral admixture of the B4-TW model, only the cementitious material content of factors B4 model .50 spread time of 16 s. The initial setting time and final setting time of the concrete were 539 min and 729 min, respectively.The SCC concrete had a slump flow value of 570 mm, and TThe average compressive strength and modulus of elasticity of SCC were tested on 12 specimens for each age at 7, 14, 28, 56 and 91 days, as shown in The CFT combines concrete and steel tubes so that the stress that impacts on these specimens is not the same as plain concrete. It is well known that creep and shrinkage of the concrete core directly affect the time-dependent deformation of CFT. If the surfaces of the components are not the same level, as presented in According to the theory of elasticity, the stress of CFT is converted to find exactly applied loading values on the specimens. The theory load of CFT can be expressed asIt is assumed that the plane remains plane under axial compression, which is the essence of strain compatibility so that the steel strain is the same as the concrete strain at all locations and can be calculated by the following equationThe load applies to the steel tube can be calculated on that of the concrete core as2); In which t′ (MPa); 2).The loading force of the concrete core of the CFT can be calculated the same as with the SCC specimens.When CFT is under axial loading, various cylinder stresses appear in the composite member, including vertical stress of the steel tube ; 2).In which 2); t′.Additionally, when the specimens are embedded with the VWSG inside the specimen, the conversion area can be obtained byThe conversion stress According to the loading force equations with the designated stress ratio (−6/MPa) depending on the applied loading, expansive agent content, and steel ratios.The amount of creep reduces on with the loading age of the concrete, as shown in When the specimens were stored in the dry environment (relative humidity of 50 ± 2%), the loading age significantly affected the creep values. The total creep compliance was 49% and 40% higher than the basic creep compliance at 375 days after loading ages of 14 days and 28 days, respectively, as observed in The elastic strain, creep strain, creep compliance and creep coefficient of the basic and total creep specimens at 375 days after loading are given in The shrinkage occurs as soon as the adsorbed water from the hydrated cement paste is removed due to the effects of cement hydration and water loss, so that the shrinkage strain was normally measured from the instant when the concrete starts to harden. The shrinkage and length change values were measured at an initial setting time, as shown in As a result, the autogenous shrinkage of CFT specimens (C-CFT-AS14) increased rapidly within 14 days after molding and the value reached approximately 71% of the shrinkage strain at 400 days. The shrinkage strain of the concrete stabilized from the third week. The issue that the confining effect of steel tube not only significantly restraints the shrinkage of the concrete but also prevents the swelling because of the expansive agent, is considered. Shrinkage of the CFT specimen increases only due to the result of the self-desiccation process as the macroscopic volume reduction in cementitious materials when the cement hydrates after the initial setting. The volume of hydration products is less than the sum of the volume of water and the volume of cement that is hydrated . In comp−6 mm/mm) at 126 days, which is equal to 44% of the SCC specimen at the same time.When the temperature is stable—the same as in the testing condition—the total strain at age t can be calculated without thermal expansion by the following equation (38)ε(t)The real creep strains can be obtained by subtracting the shrinkage strains from the measured creep strains, as illustrated in −6 mm/mm); −6 mm/mm), which is elastic strain if the stress is small; −6 mm/mm), −6 mm/mm).In which The shrinkage strain of specimens was measured after the curing of concrete in addition to the creep strain and total strain presented, and these values are presented in The measured time of shrinkage and length change tests were taken at 14 days and 28 days as well as loading age of creep tests, and the data are plotted in The strains of the CFT components were measured by separate strain gauges. As can be observed in In the initial period after loading, the steel tube had a vertical strain value higher than the concrete core because the Poisson’s ratio of concrete is lower than that of the steel. The difference of vertical strains of CFT components also indicates that longitudinal slipping between the steel tube and concrete can occur in this stage. The vertical strain of the concrete was higher with time due to creep in addition to the larger lateral strain. However, the deformation of the concrete core is restrained by the confinement effect of the steel tube so that the vertical strain of the concrete increases equally with steel. The result confirms that the plane remains plane as well as compatibility between steel and concrete components. The test data compared suitably with those of Kwon et al. , who detWhen the CFT specimens were subjected to a sustained load of 325 kN with a conversion stress ratio of 0.32, the stress of the components can be evaluated according to the data of strain gauges. −6 mm/mm) at 375 days and a modulus of elasticity of 355,555 MPa, respectively, a D/t ratio of 36, and a yield strength of steel of 250 MPa, the stress transfer of the steel tube was predicted to be 7% of steel stress equal to 17.5 MPa, which is an error of 50% in comparison with the experimental data.The stress result of the composite components varied over time and was also shown by Kwon et al.  and WangThe vertical strains, vertical stresses, and stress transfer of CFT components varied at various days, as can be seen in B4 model and B4-TW model were used to predict the creep and shrinkage of the concrete. The B4-TW model extends and refines the B4 model so that the parameter values of the two models are almost the same and are given in . HoweverZ. P. Bažant  indicateq2, q3 and q4, which are factors that affect the basic creep value, predicted by the empirical formula of the B4 model was higher than the value calculated based on the B4-TW model. Therefore, the B4 model gave much higher—from 54 to 57%—long-time values of basic creep compliance than those of the B4-TW model at 375 days after different loading ages, as illustrated in As can be seen in q5—that controls the magnitude of the drying creep compliance—calculated by B4-TW model, is 1.84 times higher than for the B4 model. Furthermore, the instantaneous strain q1, which is closely related to the elastic modulus value, estimated by the B4-TW model, is also larger than that estimated by the B4 model  and the plain concrete curing in moist conditions (C-SCC-LW14 specimen), as observed in shε∞ for the B4 model is also less than that of the B4-TW model, as shown in The total shrinkage of specimens with a curing age of 14 days (C-SCC-TS14) and 28 days (P-SCC-TS28) are compared to the values of the B4 and B4-TW prediction models, as shown in In the study, the SCC was used to fill the steel tube but the shrinkage of the SCC is higher than the plain concrete due to its higher amount of paste and lower water-to-binder ratio. The addition of expansive additive to the SCC in the CFT is common to prevent the gap caused by significant autogenous shrinkage. However, the experiment shrinkage also showed large divergences in comparison with the predicted values of the models, as presented in (1)The advantage of the interaction between the steel tube and the concrete core significantly reduced the time-dependent deformation of CFT. The steel tube has a role as a restraint on the concrete, creating an enclosed environment that prevents the concrete spalling and the relaxing of the impacted stress. Therefore, the creep compliance and shrinkage strain of the CFT specimens was33% and 62%lower than that of the concrete without the steel tube, respectively.(2)The creep strain in the moisture condition was lower than that in the drying condition. Additionally, the various loading ages that affected the total creep were higher than basic creep because the basic creep is not influenced by the level of humidity, the total creep has a rapid drying process at an earlier age. However, the different curing ages negligibly affect the total shrinkage and length change of the specimens.(3)The vertical strain of the CFT components had a similar magnitude which was only 1.1% different compared to the vertical strains of the steel tube and the concrete core so they can be considered as fully composite members.(4)The vertical stress of the steel increased by 32.7% but the vertical stress of the concrete decreased by 15.8% at 375 days. This stress transfer magnitude cannot be neglected and should be considered in CFT design.(5)The B4-TW model is more consistent than the B4 model in predicting the basic creep of the concrete. The prediction values of the two models were close to the test data of the short-term total creep under about 100 days but higher than those of the long-term total creep. The B4 model does not predict the autogenous shrinkage of concrete with slow hardening cement. Additionally, both the B4 model and the B4-TW model are not suitable to estimate the shrinkage of SCC using expansive admixture.This paper tests the creep and shrinkage behavior of CFT and concrete without steel tube specimens. The SCC used an expansive admixture to fill a steel tube of 108 mm in diameter, 3 mm in thickness, and 300 mm in length. The time-dependent deformation was performed under different curing conditions over a long time period of up to 400 days. Based on the experimental results, the main conclusions of the study can be drawn as follows:


A series of NiAl-Cr-Mo systems were produced and assessed as far as their microstructure and their sliding wear resistance is concerned. The NiAl content was kept constant and seven compositions of Cr-Mo were tested, namely, 40Cr-0Mo, 30Cr-10Mo, 25Cr-15Mo, 20Cr-20Mo, 15Cr-25Mo, 10Cr-30Mo, and 0Cr-40Mo. It was observed that most of the systems contained primary phases, eutectic microconstituents, and, occasionally, intermetallic phases as the outcome of peritectic reactions. The extent and the nature of all these microstructural features was proved to be affected by the Cr/Mo relative ratio, and an attempt was conducted in order to explain the microstructural features based on solidification and other related phenomena. It was observed that the increase of the relative Mo/Cr ratio led to a significant restriction/elimination of the eutectic microconstituent. The sliding wear response of the produced system seems to diverge from the classical sliding wear laws of Archard and is based on multiple factors such as the nature of the oxide phases being formed upon sliding, the nature and the extend of the intermetallic phases being formed upon solidification, and the integrity and rigidity of the primary phases—last to solidify areas interfacial region and the factors that may influence this integrity. During the last decades, NiAl-based alloys have been considered as potential candidates for high-temperature applications due to their enhanced properties at such extreme servicing conditions and possible substitutes for Ni-based superalloys . HoweverWithin this frame, the addition of Cr within a NiAl intermetallic matrix has gained special scientific and research interest ,8. The pIn parallel with the works focused on the addition of Cr in NiAl matrix, Mo has also attracted significant research attention, and various works have been conducted concerning the development and microstructural property evaluation of NiAl-Mo systems with very promising results ,5,17,18.By the development of HEAs point of view, a combination of Cr and Mo would be an interesting approach of gradually building medium to high-entropy systems. This combination of four elements (Ni-Al-Cr-Mo) has been examined in various research efforts. Chen et al. , for insDespite the fact that the microstructure and the mechanical properties of NiAl-Cr-Mo systems were examined through various research efforts, limited work has been conducted so far on the sliding wear response of these systems. The most systematic approach on this issue has been conducted by Guo and his colleagues  who examThe present work is part of a wider research effort to design and develop new high-entropy alloys around the fundamental core of the NiAl-Cr system. The authors have, at a first instance, investigated the basic NiAl-Cr core system, as far as the microstructural features and the related mechanisms are related, along with the alterations that small additions of Mo may cause . In the High-purity elemental powders were mixed in the intended atomic ratio in order to produce various Ni-Al-Cr-Mo alloys. Approximately, 5 g of raw materials, with purities higher than 99.5%, were subjected to uniaxial compression in order to make green pellets with targeted  composition according to α radiation (λ = 1.5406 Å) and a secondary beam graphite monochromator. Samples were scanned over an angular 2θ range from 20° to 120°, in steps of 0.02° (2θ) at a rate of 2 s per step .The microstructure of the alloys was studied in their as-cast condition. The metallographic specimens were mounted in Bakelite, abraded on SiC papers up to 2400 grit, and then polished with a 3 μm diamond suspension. The polished specimens were etched using aqua regia. Alloy’s microstructure was analyzed with the use of a scanning electron microscope (SEM)  equipped with both backscatter electron (BSE) and energy-dispersive spectroscopy (EDS) detector of X-Act type by Oxford Instruments . X-ray diffraction (XRD) patterns were obtained using a D8 ADVANCE diffractometer , operating with CuK1) was measured using a 136° diamond pyramid for 5 s of load (1 kg) application. Sliding wear tests were performed using a ball-on-disk apparatus  at ambient temperatures with relative humidity around 70%. The externally applied load was 5 N. The overall sliding distance was 1000 m, and the sliding speed was 10 cm/s. A 100Cr6, 6 mm steel ball was used as a counter body material. All specimens were acetone cleaned and their weight was measured. Tests were interrupted every 200 m, and the specimens were reweighted. Wear tracks were examined by SEM-EDS analysis. Three independent runs were carried out for each individual alloy system.Vickers microhardness , which is a strong indication for potential phase segregation in their case. In the case of the geometric factor γ proposed by Wang et al. [f < 37 meV/atom range for single-phase stability. Finally, the model of Senkov et al. [g et al.  model ong et al.  model, vg et al.  that deag et al.  model gav et al. , which aA possible concluding remark, as the parametric model predictions are concerned, is the fact that there is no consistent tendency for either a single-phase solid solution or multiple-phase formation. The actual microstructures do verify the presence of multiple phases, as it will be discussed in the following paragraph. This skepticism on the parametric model prediction validity, which has been expressed by many researchers , the microstructure is lightly modified by means that the eutectic phase is reduced and the dark interdendritic phase is expanded. EDS analysis reveals that inside the core of the primary dendrites, the Mo content is increased and the Cr content is decreased. The dendritic halo is rich in Ni, as in the previous case, and the dark phase is rich in Ni and Al. Despite the reduction of the eutectic area, its morphology is still of typical lamellar configuration.According to In the case of alloy E, where the Mo content reaches the values of 25 at.%, the microstructure is further modified. The extent of the peripheral halos is increased as the extent of the dark phase also does. It is worth noticing, however, that the eutectic morphology is almost eliminated. Further increase of the Mo content, alloy F in Alloy G (no Cr addition), 3 and MoNi). AlMo3 seems to vanish for high Mo additions (over 20%Mo), whereas MoNi persists for even higher Mo concentrations. The NiAl-40 Mo alloy contains practically Mo and NiAl phases. Another important observation is the fact that from alloy A to alloy G, there is a gradual shift of one of the BCC phases, that of Cr, from Cr in the case of alloy A to Mo in the case of alloy G, a fact that is also expected. Despite the fact that TEM examination could provide more details on the peritectic halos formation, such examination, as adopted in other experimental efforts [Apart from the various EDS analysis, XRD measurements complement the characterization of the various phases being formed in the examined alloys . It can  efforts ,28, was The first important observation to be noticed is the fact that the microstructural features of each individual alloy system produced in the present effort and presented in It was shown in Since this alloy contains no Mo, an approach to explain the various microstructural features should be passed on the partial binary eutectic system. As such, according to the NiAl-Cr pseudobinary phase diagram proposed by Tang et al.  and DuprOnce the primary B2 is formed, further heat release due to recalescence causes temperature increase at the surrounding liquid above T0-B2, establishing conditions for eutectic undercooling, leading to the formation and growth of eutectic microconstituent. The alterations of the eutectic morphology observed in this system are also attributed to recalescence phenomena: As soon as temperature rises above the T0-B2 point, the undercooling for eutectic growth is considerably high. High undercooling favors rod- or hive-like (commonly referred as fibrous) eutectic morphologies . Taking It has been already mentioned in a previous paragraph that the overall microstructure complies with the predictions suggested by Peng et al.  accordin(1)Based on EDS analysis , the pri(2)The presence of Mo most likely affects the T0-A2 temperature presented by Tang et al. . The pre(3)Synergistically to these potential mechanisms, another parameter that may also enhance the stability of the primary A2 Cr(Mo) phase is the lattice distortion value δ. As shown in This A2 primary phase stabilization can be explained by the following phenomena that can act either individually of synergistically:3 intermetallic phase. Further temperature decrease will drive the system to a second peritectic reaction to occur, that of the Mo-Ni system, forming a second zone of peritectic reaction product, that of the MoNi intermetallic phase.These potential phenomena can explain the formation of the core of A2 primary phase. However, as indicated in It is proposed thus for the primary phase that after the initial formation of the basic Cr(Mo) core, a sequence two zones of peritectic reaction zones is following, leading to the formation of the characteristic halo configuration peripherically to the basic Cr(Mo) core.The solidification process is concluded with the formation of the eutectic microconstituent. EDS mapping analysis  on the eThe general view of this alloy system, as presented in 3 intermetallic phase). Once the first peritectic zone is formed, further temperature reduction leads to the second peritectic reaction and the formation of the second peritectic zone product, that of the MoNi intermetallic.As far as the primary phase is concerned, no significant alterations of its formation sequence and mechanism can be noticed. Mo content in the primary phase core is raised  and in conjunction with the further increase of the liquids’ temperature and the T0-A2 temperature, ensures the stability and the domination of the A2 phase. As far as the halo phase is concerned, the Mo/Al ratio in the center of the core is roughly 90/10, and as such, the formation of a primary MoAl peritectic reaction product is not expected. Indeed, EDS analysis  did not Questions are arisen, yet, as far as the solidification of the last liquid is concerned. EDS mapping analysis of this area revealed a further decrease of the Cr content close to 15 at.%. The obtained microstructure suggests that the depletion of Cr by the primary phase shifts the composition of the remaining liquid even deeper to the hypoeutectic area, and following the behavior Mo may have on the liquidus and the T0-B2 temperatures, it is most likely that the experienced undercooling provided the conditions for primary B2 phase formation, yet recalescence and heat release did not provide extended undercooling conditions for eutectic constituent formation. As such, the eutectic area is further suppressed.The overall microstructure is in compliance with the predictions of Peng et al. . In-deptAnother important observation , howeverThis final alloy contains no Cr, and as such, its microstructural features should be approached by the pseudobinary NiAl-Mo phase diagram, as proposed by Peng et al. . Based o(1)Upon solidification initiation, undercooling falls below the T0-A2 (Mo) and T0-B2 (NiAl) temperatures establishing such conditions for decoupled, partitionless growth of Mo and NiAl phases, with Mo being the first one to form due to its higher melting point.(2)As soon as recalescence is expressed, temperature is raised so that Mo continues to grow in a partitioning to the liquid mode, whereas NiAl proceeds in a partitionless way.(3)Further increase of temperature in the remaining liquid establishes conditions for partitioning growth of both NiAl and Mo establishing conditions for the development of the eutectic mixture growth, as the final stage of the solidification process.Under this phenomenologically logical assumption, the solidification sequence can be as follows:Another important observation is the lack of halo zones at the periphery of the light phase, which is an indication of significant suppression of the peritectic reaction product. Indeed, as EDS analysis indicates , the relWhen focusing in (1)A peak value of the wear rate is observed for alloy B, which is gradually reduced until the case of alloy F. This practically means that the introduction of small amount of Mo bursts the wear rate at very high values, yet the progressive increase of the Mo content leads to a gradual reduction of the wear rates.(2)All wear tracks  show an (3)The oxide phases being present within the cavities of the removed materials are also of great importance. EDS analysis  on these(4)Even among these alloys, a classification of their wear rates can be conducted: alloys B and C forming the first group of high wear rates, alloys D and E forming a second group with reduced wear rates, and finally alloy F with considerably reduced wear rate. There must be a reason for this kind of transition which will be addressed later.(5)Another observation that will help in the explanation of the wear behavior of the different systems is the morphology of the primary phases and, especially, the morphology of the contour of the primary phases. By recalling the involved microstructures , it can The introduction of Mo, resulted in the expression of an obscure situation: despite the fact that the presence of Mo leads to an increase of microhardness, the wear resistance does not seem to follow the same trend. A close examination of the wear track morphologies of these systems  reveals (1)The first crucial factor, as already mentioned, is the nature of the oxides formed. The observed oxides in these systems are Fe-based, Ni-based, Mo-based, and Cr-based oxides. According to the extensive reviews of Glascott et al.  and Stot(2)The severe mass loss, as already mentioned, is related to the extensive detachment of the last to solidify liquid areas. It should, hence, be a gradual weakening of the interfacial area between the primary phases and these lastly solidified regions. The difference between alloy A and alloys B to F, concerning these interfacial areas, is the presence of peritectic halos around the primary phases. As mentioned previously, the result of these peritectic reactions is the formation of strong and brittle intermetallic phases. Taking into consideration that as the Mo content increases from alloys B to C, the second primary phase to be formed is the NiAl intermetallic phase; Based on these important observations, a possible explanation of this category of system wear response could be as follows:The presence of intermetallic phases perimetrically to the primary grains, can also affect the integrity of the interfacial areas in other terms. As mentioned previously, the contour of the primary phases in alloys B to F has a strong acicular jigsaw morphology. This aggressive landscape can also establish a field of high residual stresses that may lead to a premature collapsing of the interfacial area. The acicular morphologies are also the result of the formation of intermetallic phases, which according to the classic nucleation theory shows increased values of the so-called Jackson factor , which is an indication of a faceted growth of mode that, in turn, leads to the formation of such kind of acicular morphologies .Based on the remarks mentioned for the alloys in the previous case, in this alloy, the dominant oxides are Ni, Fe, and Mo , i.e., oIn summary, the wear response of the alloys produced on the present effort is a combination of the factors reported previously, i.e., the nature and chemistry of the surface oxides and the presence, the nature, and the morphological features of intermetallic phases. It has to mentioned, nevertheless, that examination of cross section areas related to the wear tracks, as adopted in other experimental efforts ,63, woul(1)The microstructure of the NiAl-40Cr alloy is a result of solidification sequence according to which primary A2(Cr) and B2(NiAl) phases are developed partitionless due to undercooling phenomenon. Recalescence and its involved heat release lead to the development of a final eutectic structure.(2)3 and MoNi intermetallic phases are developed at the periphery of the primary grains due to peritectic reactions. Recalescence leads to the final stage of eutectic phase formation.The substitution of Cr by 10% Mo, in the case of alloy B, stabilizes the A2 phase as a primary phase as both Mo and undercooling do not favor a parallel B2 growth in a partitionless way. Phase diagram predictions verified in practice and multiple AlMo(3)3 is eventually vanished, and the eutectic phase is continuously restricted. B2 (NiAl) is gradually increased as a secondary (following A2) “primary” phase. All these microstructural sequences are related to partitioning/partitionless growth, recalescence, and undercooling phenomenon.By increasing the Mo content in all the preceding alloys, the A2 phase is further established as the primary phase, AlMo(4)The NiAl-40Mo alloy consists of primary Mo, NiAl, and NiAl-Mo eutectic constituents. The same phenomena, as in the previous cases, are responsible for this solidification sequence.(5)The increase of the Mo content leads to progressive elimination of the eutectic microconstituent and establishes significantly simpler microstructures of greater controlling potential.(1)The sliding wear behavior of alloy NiAl-40Cr showed extensive material loss due to localized delamination of the eutectic areas. Oxides of Fe and Cr are the main oxides being present.(2)An increased wear rate was observed in the case of alloy B (NiAl-30Cr-10Mo) by extensive material loss at the eutectic areas, due to the presence of intermetallic peritectic halos and the nonlubricating action of the formed oxides. The increase of the Mo content leads to the reduction of the wear rate , due to the gradual restriction of the intermetallic phases and the gradual formation of oxides phases with intensive lubrication action.(3)The NiAl-40Mo is characterized by a relatively low wear rate due to the absence of intermetallic phases and the presence of lubricating oxides.