Introduction

Result and discussion

Conclusion

Acknowledgment

Appendix A. Supplementary data

Research Data

References

Abstract

In this paper, the effect of Zr and Mn on the microstructure and first hydrogenation kinetic of TiFe alloy is reported. TiFe alloy to which Zr, Mn or a combination of both have been added were synthesized by induction melting. First hydrogenation of all alloys was performed at room temperature under 20 bar of hydrogen. We found that addition of manganese makes possible activation at room temperature, but kinetics was very sluggish. Alloy with 2 wt% Zr did not absorb hydrogen. However, with addition of 4 wt% Zr, the alloy absorbed 1.2 wt% of hydrogen. A synergetic effect was found when zirconium was added along with manganese. Alloy with 1 wt% Mn and 2 wt% Zr had better kinetics than the alloy having only Mn or only Zr. The maximum hydrogen capacity was also greater at ~1.8 wt% after 7 h. Combination of 4 wt% Zr and 2 wt% Mn absorbed 2 wt% of hydrogen in 5 h. The rate limiting step for each activated alloy was found to be diffusion controlled with decreasing interface velocity.

Introduction

Emission of greenhouse gases due to the intensive use of fossils fuel demands the development of alternative fuels [1,2]. Hydrogen fulfills the criteria as an alternative energy carrier due to its high energy density, availability and less impact on the environment when produced by renewable resources such as solar and wind energies [3,4]. One of the major challenges in the development of hydrogen economy is to store hydrogen safely and at low cost [5]. The solid-state hydrogen storage in metal hydrides presents some advantages over the conventional high-pressure cylinders and liquid hydrogen because of its high volumetric capacity, low pressure, and low temperature of operation [6e10]. Metal hydrides can also be used as a negative electrode in rechargeable batteries such as Ni-MH [11e13]. Two of the major characteristics of metal hydrides should have in order to fulfill the requirements for mobile and stationary applications are low cost and utilization in a practical range of temperature and pressure (0e100 C, 1e10 atm) [14e16].

1- Introduction

2- Materials and experiments

3- Model formulation

4- Model implementation

5- Numerical results

6- Conclusions

References

Abstract

The one-way and two-way shape memory effects (SMEs) as well as the thermal hysteresis represent fundamental properties when dealing with the design of detachable and thermally-stable connection systems based on shape memory alloys (SMAs). Such properties can be induced and tuned by thermo-mechanical processes that include thermal treatments and severe pre-deformation in martensitic state, causing the onset of plastic strains. In such complex conditions, material modeling is of great importance to support the design. This paper proposes a generalization of the three-dimensional phenomenological constitutive model by Souza et al. (1998), in order to describe the behavior of severely pre-strained NiTi-based SMAs. The proposed model allows to describe pseudoelasticity, one-way and two-way SMEs, as well as additional physical phenomena evidenced experimentally, such as transformation temperatures’ evolution, thermal hysteresis, phase transformations at low stresses, thermal strains, and phase-dependent elastic properties. Several numerical simulations, ranging from uniaxial tests to the finite element analysis of two case-studies, are performed. Model results are in good agreement with the results of a performed experimental campaign and allow to discuss SMA behavior under such complex loading conditions.

Introduction

Shape Memory Alloys (SMAs) are widely applied in various fields of engineering and medicine thanks to their unique strain recovery capabilities (Jani et al., 2014; Otuska and Wayman, 1998; Duerig et al., 1990), that can be obtained either by heating or by stress release, through the so-called Shape Memory Effect (SME) and Pseudoleastic Effect (PE), respectively. SME and PE are due to reversible solid-state microstructural transitions, the so-called thermoelastic martensitic transformations (TMT), between the parent body-centered cubic austenitic phase (B2) and the product monoclinic martensitic one (B19’). TMT can be induced either by temperature changes (TIM, ThermallyInduced Martensite) or by mechanical loads (SIM, Stress-Induced Martensite), providing the SME and PE, respectively. B2 phase is stable at high temperature and low stress, while B19’ is stable at low temperature and high stress. Thanks to TIM, the material is able to change its microstructure reversibly when varying the temperature between the so-called Transformation Temperatures (TTs). Direct transformation (B2-B19’) can occur when cooling the material below the martensite TTs, while reverse transformation takes place during heating above the austenite TTs. The difference between martensite and austenite TTs, namely ΔTT, is regarded as the thermal hysteresis of the material. Thermal hysteresis, and the associated stress-strain hysteresis, is a limiting phenomenon in SMA-based fast actuation systems, as it represents a source of inefficiencies due to the increased energy loss and time response. Nevertheless, thermal hysteresis could be a beneficial feature when a large thermal stability range is needed, such as in some fasteners or connectors operating in constrained recovery mode (Jani et al., 2014; Duerig et al., 1990; Humbeeck, 2001). In such systems recovery forces are generated as the SMA element, previously deformed in martensitic state, is heated up to the austenite TTs in a constrained shape, i.e. against a mechanical obstacle. This is due to the so-called One-Way Shape Memory Effect (OW-SME) and it is the basic principle of many successful systems, e.g. permanent fasteners, couplers, or SMA hybrid composites (Humbeeck, 2001; Duerig et al., 1990; Jani et al., 2014; Kirkby et al., 2008).

Flexibility in metal forming is needed more than ever before due to rapidly changing customer demands. It paves the way for a better control of uncertainties in development and application of metal forming processes. Although flexibility has been pursued from various viewpoints in terms of machines, material, process, working environment and properties, etc., a thorough study of the concept was undertaken in order to with problems of manufacturing competiveness and tackle new challenges of manufacturing surroundings. Therefore, in this paper, flexibility in forming is reviewed from the viewpoints of process, material, manufacturing environment, new process combinations and machine–system–software interactions.

Introduction

Background and motivation

The manufacturing sector undergoes drastic changes in the global scale due to the advent of a new industrial revolution, the so-called fourth industrial revolution or Industry 4.0 [170]. A key feature of this revolution is the increasing interaction of human beings, production machines and products. As a consequence, conventional mass production is changing gradually towards manufacturing systems for mass customization, with higher flexibility and on-demand manufacturing, which brings about new technological challenges. Such a new industrial wave demands connectivity, automation and intelligent systems, in mass customization and on-demand manufacturing in addition to mass production as summarized in Fig. 1 [184]. The mentioned connectivity and intelligentization enable realization of automated manufacturing of products and related automated services. The pursuit of these three factors would finally lead to super-connection and super-intelligentization of human beings, things and space, and ultimately resulting in systemic innovation of industry and society. While flexibility is inherent to manufacturing processes such as machining and welding, die-based manufacturing processes such as metal forming and casting are increasingly challenged to meet the demands on flexibility. A working group study on mega-trends and the future of metal forming [127] in 2012 revealed the importance of flexibility. Theresults of the study were summarized as shown in Fig. 2. Various areas, including flexibilitization, intelligent processes, complexity of shape and integration of functions, all of which demand for flexibility of forming either directly or indirectly. 3D printing of metals was included as a separate category, though not having been included in traditional forming, but it becomes a very powerful means of highly flexible manufacturing for producing metallic parts and dies with reasonable productivity for some industrial manufacturing sectors. Flexibility of metal forming has become an important issue in recent manufacturing demands and plays an essential role in customized mass and individualized production in the manufacturing industry. As flexible sheet metal forming, single point incremental forming has been widely employed and subject to extensive research [61,73]. The introduction of double point (sided) incremental forming had enhanced the formability and extended the range of shape complexity [125,199]. The combination of incremental forming and conventional sheet forming processes further extended the capability of incremental forming with additional advantages including increased productivity [13]. Incremental bulk metal forming has a long history since the Bronze age. Since then, incremental bulk forming has been developed by employing various forming processes including forging typically and many incremental forming processes using rotating tools such as flow forming, rotary swaging, orbital forming, gear rolling, pierce rolling and ring rolling, etc. [52]. Various flexible forming processes for both sheet and bulk can be found in Japan and they are well summarized in Ref. [8]. Incremental forming is surely the key flexible forming process and has been diversified for manufacturing various metallic products. However, flexibility in forming has been pursued from diverse viewpoints in terms of complexity, material, process, working environment and machine, etc. In this work, flexibility is newly viewed and subjected to review from various angles of perspective in order to provide new prospects of flexibility in metal forming.

Conclusions

References

The aim of this study is to evaluate a full integrated modelling strategy to evaluate the influence of casting defects on the fatigue life directly from process simulation. We have shown that defects characterized by their size and the microstructure characterized by the SDAS, are the main parameters that control the fatigue limit. A fatigue criterion that already takes into account for the effect of the defect on the fatigue limit was modified to introduce the effect of SDAS. This improved criterion has been employed to predict the Kitagawa diagram for multiaxial loading for different loading cases. The simulation of the modified criterion showed that the reduction of the fatigue limit with the defect size and SDAS is well described. In the last part a numerical model was developed to perform a simulation of the fatigue limit starting from the simulation of the casting process. Using this numerical model, we simulated the defect size and SDAS depending on the solidification time, the fatigue limit is simulated using the improved criterion. We proposed in this part a mold which let to obtain samples with two different microstructures. In this study, a second fatigue tests was carried out on these samples to validate the numerical simulation on the proposed mold. It turns out that the numerical model provides reasonably well the obtained experimental results.

HCF studies on cast aluminum alloys have shown [1, 2] that the effect of defect size is in competition with other microstructural features such as SDAS/DAS to limit fatigue life. It has been clearly observed that casting defects have a detrimental effect on fatigue limit above a critical size [3, 4]. In addition, the influence of microstructure defined by DAS or SDAS on the fatigue behaviour of cast aluminum alloy cannot be neglected [5-7]. It has been demonstrated [6] that the eutectic phase can cause micro-cracking due to strain accumulation in proximity to Siparticles. As reported in [1], for small defect sizes, there is a prominent interaction between defects and microstructure on the fatigue limit of cast A356-T6 aluminum alloy. Therefore, the simulation of defect size and microstructure distributions within a cast aluminum alloy component are critical inputs for fatigue design. In this context, several studies [8-12] have proposed models to simulate the solidification process and the formation of porosity. A number of models combine the shrinkage pressure and hydrogen concentration to describe porosity formation. In these models the defects were explicitly assumed to be spherical [13-15]. Atwood et al [11] assumed that during solidification, due to solid-liquid balance, pore shapes become complex as they grow between grains or between dendrite arms. Several studies [8, 13] have suggested that it is more appropriate to consider an equivalent pore size rather than trying to characterize the complex 3D shape. This assumption is useful in simplifying the fatigue design for casting aluminum alloy components. In some studies [8, 10-12], the effects of pressure, cooling rate and the distribution of hydrogen in the casting on the formation of pores has been studied.

Abstract

This study investigated the corrosion and stress corrosion cracking susceptibility of 316L stainless steel exposed to 550 °C and 600 °C supercritical water. XRD, SEM, EDS and EBSD analyses have been carried out to characterize the microstructure of the tested specimens. Results indicate that increasing the test temperature reduced the corrosion and SCC resistance of 316L SS. The oxide scale of corroded specimen shows a duplex structure, and the fracture analysis indicates that the formation of Fe3O4 ‘channels’ in the Fe-Cr spinel layer contributes to the intergranular and transgranular SCC of the materials in 550 °C and 600 °C SCW.

Introduction

Generation IV reactors are attractive due to their high efficiency and safety [1,2]. The supercritical water cooled reactor (SCWR) is one of the most promising of the six innovative Generation IV reactor concepts because it offers advantages, such as simplification of components and higher efficiency [3,4]. By operating above the thermodynamic critical point of water (374 °C, 22.1 MPa), a single phase coolant is used in SCWR and many components, such as steam generators, steam separators, dryers, recirculation and jet pumps are not necessary [5,6]. Thus, the design of the reactor is greatly simplified.

In SCWR, the fuel cladding experiences the most severe conditions among the in-core structural components. During normal operation, the temperature on the fuel cladding can be 600 °C, and it can be even higher during transients. The higher the operating temperature, the higher the possibility that the fuel cladding material will fail due to creep [7], corrosion [6,8] or stress corrosion cracking (SCC) [6], which can result in leakage of fission products and pose great challenges to the reactor safety.

Stainless steels (SS) [8,9], nickel based alloys [10,11] and oxide dispersion strengthened steels (ODS) [12,13] have been proposed as candidate materials for fuel cladding and other core components in SCWRs, and extensive research has been conducted to examine the corrosion and SCC performance of those materials [14–17]. Among the candidate materials, austenitic stainless steels (SSs) are the least expensive, and some research [6] indicates that their corrosion resistance are better than nickel-based alloys. However, other research [18,19] reported that the corrosion rate of 316L SS was much higher than nickel-based alloys.

Abstract

In this study, the hot extrusion process was applied to stir cast aluminum matrix–SiC composites in order to improve their microstructure and reduce cast part defects. SiC particles were ball milled with Cr, Cu, and Ti as three forms of carrier agents to improve SiC incorporation. Large brittle ceramic particles (average particle size: 80 μm) were fragmented during ball-milling to form nanoparticles in order to reduce the cost of composite manufacturing. The experimental results indicate that full conversion of coarse micron sized to nanoparticles, even after 36 h of ball milling, was not possible. Multi modal SiC particle size distributions which included SiC nanoparticles were produced after the milling process, leading to the incorporation of a size range of SiC particle sizes from about 50 nm to larger than 10 μm, into the molten A356 aluminum alloy. The particle size of the milled powders and the amount of released heat from the reaction between the carrier agent and molten aluminum are inferred as two crucial factors that affect the resultant part tensile properties and microhardness.

Introduction

The use of Al–Si (in particular A356) alloys in themanufacture of automotive engine components has increased considerably in recent years [1–6]. Although high-strength aluminum alloys have been developed, the addition of alloying elements and microstructural modification, can be costly, contain toxic elements, and often results in properties, which only result in a slight increase in stiffness. The demands for lightweight, high-modulus, and high-strength materials have therefore led to the development of aluminum matrix nanocomposites (AMNCs) [2,7–18].

Abstract

It is often observed that the plastic deformation of crystalline polymers, especially on drawing, causes significant amount of cavitation. Cavitation can be suppressed in the process of rolling due to compressive component. An innovative method of obtaining highly oriented polymeric materials is by unidirectional rolling in a channel formed on the circumference of one roll with another roll, having the thickness matching the width of the channel. The side-walls of the channel on one roll constitute the side constraints while the other roll works like a plunger in channel-die. Rolling with side constraints is cavity-free and is advantageous over conventional rolling, channel-die compression or solid state extrusion since it gives the possibility of obtaining relatively thick and infinitely long highly-oriented shapes or profiles in a fast continuous manner. The resulting profiles can have considerably high cross-section area. The design of the rolling machine is described. Few examples of rolling with the rate 4.23 m min−1 of isotactic polypropylene (iPP) and high density polyethylene (HDPE) shapes are presented. Tensile strength of the rods rolled to the compression ratio 5.4–6.6 with final cross-section of 10–12 mm approached 200 MPa for both polymers. The oriented rods of iPP and HDPE demonstrated high and sharp texture produced in HDPE by the activity of (100) [001], (010) [001] and (100) [010] crystallographic slips while in iPP (010) [001], (110) [001] and (100) [001] slip systems were active. In addition to those slips twinning modes were active on unloading. Appropriate combination of rolling rate, temperature, initial thickness of rolled bar, as well as molecular weight of the polymer will apparently lead to rods with higher strength.

Abstract

In the present work, SiCp/2024 composite powders were produced by low energy ball milling a mixture of centrifugally atomized 2024 particles and β-SiC reinforcements having a size range between 2 and 6 μm. The experimental parameters studied were the milling time (between 1 and 24 h) and the volume proportion of β-SiC in the starting powder mixture (between 5 and 35%). These SiCp/2024 composite powders were consolidated by hot extrusion using optimized processing parameters. The microstructure of the extruded composites showed a uniform distribution of β-SiC particles in the 2024 aluminum matrix. Moreover, the mechanical properties (yield strength and Young’s modulus) improved with the SiC content although some reduction in ductility was observed.

Introduction

For more than two decades, aluminum-matrix composites have been made by different methods including casting [1–9], spray atomization and codeposition [9,10], pressure-infiltration of porous ceramic preforms [11], and powder metallurgy techniques like densification of mixtures of metal and ceramic powders [12–17], or of composite powders made by mechanical processing [13,18–34]. The mechanical properties of these composites are influenced significantly by the quantity and mode of distribution of the reinforcements and also by the nature of the interfaces between the reinforcements and the matrix.

Shape memory alloys (SMAs) are an exquisite class of active materials with an ability to regain its original shape at high temperatures. There is a wide range of alloys which exhibit the shape memory effect, but only those alloys are commercially attractive which show a substantial amount of strain recovery and generate significant force due to shape change. Among the Cu-based shape memory alloys, Cu-Al-Ni alloys have a higher thermal stability than Cu-Zn-Al alloys [1, 2]. Therefore, Cu-Al-Ni alloys are being developed for high temperature applications due to their potential to be used as sensors and actuators at temperatures around 200 ºC. On the other hand, Cu-Zn-Al alloys have maximum working temperatures of 120 °C, but they sh ow better ductility as compared to Cu-Al-Ni alloys for low temperature applications [2]. Shape Memory Effect (SME) is shown by alloys exhibiting crystallographically reversible thermo-elastic martensitic transformation. At a higher temperature, the same alloys have another unique property called as super-elasticity [3]. Superelasticity is caused due to large non-linear recoverable strain (up to 18%) upon loading and unloading, in which a specimen once deformed by application of force regain its original shape automatically without any application of heat [1]. The Ni-Ti, ferrous alloys and Cu-based alloys are considered as practical materials for applications among many shape memory alloys. Ni-Ti alloys, an equi-atomic compound of Ni and Ti, are the most widely used shape memory alloys. They show excellent shape memory strain up to 8 % and are thermally stable. However, the reactivity of Ti limits their processing in air and hence all melting operations are to be carried out in vacuum. In recent decades, Cu-based shape memory alloys have emerged as a potential material for variety of applications, such as high damping material, sensors and actuators. Cu-Al-Ni shape memory alloys have gained special attention due to their high thermal stability among the other Cubased shape memory alloys. The presence of SME, thermo-elastic martensitic transformation and crystallography in the Cu-Al-Ni alloy was confirmed by Otsuka [4-6]. Some ferrous alloys also exhibit SME under certain conditions [7]. Fe-Mn-Si alloys are the most important iron-based shape memory alloy. However, they can recover shape memory strain less than 4 %.

Shape memory alloys have the ability to undergo a reversible thermoelastic phase transformation from the parent phase austenite to a product phase martensite. This particular behaviour is responsible for two main functional properties: the superelasticity and the shape memory effect [1]. In the former one, the martensitic transformation is induced by applied thermomechanical loading of the austenitic phase, leading to important inelastic strains which are recovered upon unloading due to the reverse phase transformation to austenite. Austenite can transform into martensite during mechanical loading, leading to an inelastic deformation that can be recovered spontaneously upon unloading [2]. For the shape memory effect, the material is deformed in the martensitic state, leading to permanent inelastic strains upon unloading. These can be removed by heating up the material to the austenitic phase again: the initial shape is restored triggering the shape memory effect [3]. Cu-based shape memory alloys are considered to be possible substitutes of NiTi, the most used shape memory alloy, owing to their lower production costs and recent improvements on their mechanical properties [4,5]. There are different Cu-based systems, with the most important ones being Cu-Al-Mn [6], Cu-Al-Ni [7] and Cu-Al-Be [8]. The latter one presents, aside from the superelastic and shape memory effects, other interesting properties: strong sound, vibration and mechanical damping capability; high mechanical strength; and resistance to corrosion [9]. These alloys exhibit also a particular technological interest for low and intermediate temperature applications [8]: as a result of the introduction of small amounts of Be, the transformation temperatures decrease drastically, allowing for the occurrence of superelasticity at very low temperatures. In fact, some empirical formulae have been proposed to determine the transformation temperatures in Cu-Al-Be alloys based on their composition [10]. Thus, these alloys can be used in a wide range of temperatures, at ambient temperature or above, but they can be considered also for cryogenic applications [9]. One of the potential applications of Cu-Al-Be alloys is in seismic devices [11] as they combine several properties: high fatigue resistance, coupled with a low sensitivity for changes in the mechanical behaviour for near-ambient temperatures; they are not very sensitive to frequency changes in the 0.1 to 5 Hz regime. Moreover their properties suffer slow degradation due to the surrounding environment [12].

مقاله سال 2016

در این مطالعه، کامپوزیت های ماتریس فلزی آلومینیوم تقویت شده با نانوذرات ZrO2 در کسر حجمی 0.5، 0.75 و 1 درصد از طریق فرایند (ARB) اتصال نورد تجمعی تولید می شود. نتایج حاصله از ریزساختار نشان دهنده توزیع ذره ZrO2 عالی در ماتریس آلومینیوم پس از 10 سیکل از فرایند ARB است. همچنین نتایج انکسار اشعه ایکس نشان می دهد که کامپوزیت نانوذرات Al/ZrO2 با متوسط اندازه کریستالیته  از 48.6 نانومتر بطور موفقیت آمیزی پس از 10 سیکل از فرایند ARB به دست می آید. تست های کشش در نوارهای ARBed انجام می شوند. استحکام کششی 2.15 بار بیشتر از مقدار اولیه افزایش می یابد. کشیدگی بطور ناگهانی در سیکل نخست رها می شود، و سپس کمی افزایش می یابد. مشاهدات تصاویر SEM از سطح شکستگی نشان می دهد که پس از 10 سیکل از فرایند ARB شکست تقریبا حالت شکست را با منافذ ریز و کشیده شده برش می دهد.

درستی شکل یک جز ماشین کاری شده، یکی از شاخصهای عملکرد یک فرایند ماشین کاری محسوب میشود. میکرو ماشین کاری آلتراسونیک، یک چنین فرایندی است که در آن درستی شکل جزء میکرو ماشین شده، به طور قابل توجهی به پایداری شکل ابزار بستگی دارد. بر خلاف ماکرو ماشین کاری (ماشین کاری درشت)، یک میزان بسیار کوچک از فرسایش ابزار در میکرو ماشین کاری (ماشین کاری ریز) میتواند منجر به تغییرات قابل توجهی در درستی شکل جزء ماشین شده شود. انتخاب مناسب ماده ی ابزار برای غلبه بر این مشکل، ضروری است. مطالعه ی کنونی، اثر ماده ی ابزار، اندازه سنباده و تغذیه مرحله ای در ساخت میکرو کانالها با ماشین کاری آلتراسونیک روی شیشه بوروسیلیکات را مورد بحث قرار میدهد. توسعه ی میکرو کانالها با استفاده از میکرو ماشین کاری آلتراسونیک به ندرت گزارش شده بود. مشاهده شد ابزار تنگستن کربید یک درستی شکل بهتر را در مقایسه با میکرو کانال ماشین شده با ابزار فولاد ضد زنگ، ارائه کرد. مکانیسم فرسایش ابزار در هر دو ما با بررسی اسکن میکرو گرافهای الکترونی ابزار به عنوان مدرک، پیشنهاد میشود. یک عامل در یک دستاورد زمانی برای مطالعه ی اثر پارامترهای فرایند مختلف، بکار رفت.