Chemical composition of the 354 alloys used in this study.
\\n\\n
More than half of the publishers listed alongside IntechOpen (18 out of 30) are Social Science and Humanities publishers. IntechOpen is an exception to this as a leader in not only Open Access content but Open Access content across all scientific disciplines, including Physical Sciences, Engineering and Technology, Health Sciences, Life Science, and Social Sciences and Humanities.
\\n\\nOur breakdown of titles published demonstrates this with 47% PET, 31% HS, 18% LS, and 4% SSH books published.
\\n\\n“Even though ItechOpen has shown the potential of sci-tech books using an OA approach,” other publishers “have shown little interest in OA books.”
\\n\\nAdditionally, each book published by IntechOpen contains original content and research findings.
\\n\\nWe are honored to be among such prestigious publishers and we hope to continue to spearhead that growth in our quest to promote Open Access as a true pioneer in OA book publishing.
\\n\\n\\n\\n
\\n"}]',published:!0,mainMedia:null},components:[{type:"htmlEditorComponent",content:'
Simba Information has released its Open Access Book Publishing 2020 - 2024 report and has again identified IntechOpen as the world’s largest Open Access book publisher by title count.
\n\nSimba Information is a leading provider for market intelligence and forecasts in the media and publishing industry. The report, published every year, provides an overview and financial outlook for the global professional e-book publishing market.
\n\nIntechOpen, De Gruyter, and Frontiers are the largest OA book publishers by title count, with IntechOpen coming in at first place with 5,101 OA books published, a good 1,782 titles ahead of the nearest competitor.
\n\nSince the first Open Access Book Publishing report published in 2016, IntechOpen has held the top stop each year.
\n\n\n\nMore than half of the publishers listed alongside IntechOpen (18 out of 30) are Social Science and Humanities publishers. IntechOpen is an exception to this as a leader in not only Open Access content but Open Access content across all scientific disciplines, including Physical Sciences, Engineering and Technology, Health Sciences, Life Science, and Social Sciences and Humanities.
\n\nOur breakdown of titles published demonstrates this with 47% PET, 31% HS, 18% LS, and 4% SSH books published.
\n\n“Even though ItechOpen has shown the potential of sci-tech books using an OA approach,” other publishers “have shown little interest in OA books.”
\n\nAdditionally, each book published by IntechOpen contains original content and research findings.
\n\nWe are honored to be among such prestigious publishers and we hope to continue to spearhead that growth in our quest to promote Open Access as a true pioneer in OA book publishing.
\n\n\n\n
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Trainees as well as established consultant thoracic surgeons, anesthetists, pulmonologists, pediatricians, obstetricians, and intensivists should find this book both interesting and provocative.",isbn:"978-1-83968-066-3",printIsbn:"978-1-83968-065-6",pdfIsbn:"978-1-83968-067-0",doi:"10.5772/intechopen.73885",price:100,priceEur:109,priceUsd:129,slug:"pneumothorax",numberOfPages:86,isOpenForSubmission:!1,hash:"0b1fdb8bb0448f48c2f234753898f3f8",bookSignature:"Khalid Amer",publishedDate:"December 11th 2019",coverURL:"https://cdn.intechopen.com/books/images_new/7093.jpg",keywords:null,numberOfDownloads:2538,numberOfWosCitations:0,numberOfCrossrefCitations:1,numberOfDimensionsCitations:1,numberOfTotalCitations:2,isAvailableForWebshopOrdering:!0,dateEndFirstStepPublish:"July 9th 2018",dateEndSecondStepPublish:"July 30th 2018",dateEndThirdStepPublish:"September 28th 2018",dateEndFourthStepPublish:"December 17th 2018",dateEndFifthStepPublish:"February 15th 2019",remainingDaysToSecondStep:"2 years",secondStepPassed:!0,currentStepOfPublishingProcess:5,editedByType:"Edited by",kuFlag:!1,biosketch:null,coeditorOneBiosketch:null,coeditorTwoBiosketch:null,coeditorThreeBiosketch:null,coeditorFourBiosketch:null,coeditorFiveBiosketch:null,editors:[{id:"63412",title:"Dr.",name:"Khalid",middleName:null,surname:"Amer",slug:"khalid-amer",fullName:"Khalid Amer",profilePictureURL:"https://mts.intechopen.com/storage/users/63412/images/system/63412.jpg",biography:"Qualified from the University of Khartoum - Sudan. 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I have special interest in minimal access thoracic surgery (VATS) for lung cancer (lobectomy), thymectomy for myasthenia, mediastinal cysts and mediastinal tumours.",institutionString:"University Hospital Southampton",position:null,outsideEditionCount:0,totalCites:0,totalAuthoredChapters:"3",totalChapterViews:"0",totalEditedBooks:"1",institution:{name:"Southampton Hospital",institutionURL:null,country:{name:"United States of America"}}}],coeditorOne:null,coeditorTwo:null,coeditorThree:null,coeditorFour:null,coeditorFive:null,topics:[{id:"1047",title:"Pulmonology",slug:"pulmonology"}],chapters:[{id:"68669",title:"Indications of Surgery in Pneumothorax",slug:"indications-of-surgery-in-pneumothorax",totalDownloads:365,totalCrossrefCites:0,authors:[{id:"63374",title:"Prof.",name:"Hany",surname:"Elsayed",slug:"hany-elsayed",fullName:"Hany Elsayed"}]},{id:"65152",title:"Primary Spontaneous Pneumothorax, a Clinical Challenge",slug:"primary-spontaneous-pneumothorax-a-clinical-challenge",totalDownloads:812,totalCrossrefCites:0,authors:[{id:"188071",title:"Dr.",name:"Fabian",surname:"Giraldo",slug:"fabian-giraldo",fullName:"Fabian Giraldo"},{id:"268282",title:"Dr.",name:"Ruby",surname:"Romero",slug:"ruby-romero",fullName:"Ruby Romero"},{id:"278567",title:"Dr.",name:"Melissa",surname:"Mejia",slug:"melissa-mejia",fullName:"Melissa Mejia"},{id:"278568",title:"Dr.",name:"Estefania",surname:"Quijano",slug:"estefania-quijano",fullName:"Estefania Quijano"}]},{id:"65217",title:"Video-Assisted Thoracoscopy in the Management of Primary and Secondary Pneumothorax",slug:"video-assisted-thoracoscopy-in-the-management-of-primary-and-secondary-pneumothorax",totalDownloads:391,totalCrossrefCites:0,authors:[{id:"268004",title:"Prof.",name:"Wickii",surname:"Vigneswaran",slug:"wickii-vigneswaran",fullName:"Wickii Vigneswaran"},{id:"268007",title:"Dr.",name:"John",surname:"Costello",slug:"john-costello",fullName:"John Costello"},{id:"281023",title:"Dr.",name:"Kostantinos",surname:"Poulikidis",slug:"kostantinos-poulikidis",fullName:"Kostantinos Poulikidis"},{id:"281024",title:"Dr.",name:"Lee",surname:"Gerson",slug:"lee-gerson",fullName:"Lee Gerson"}]},{id:"65079",title:"Catamenial Pneumothorax",slug:"catamenial-pneumothorax",totalDownloads:493,totalCrossrefCites:1,authors:[{id:"268979",title:"Prof.",name:"Sezai",surname:"Celik",slug:"sezai-celik",fullName:"Sezai Celik"},{id:"279787",title:"Dr.",name:"Ezel",surname:"Erşen",slug:"ezel-ersen",fullName:"Ezel Erşen"}]},{id:"68427",title:"Controversies in Pneumothorax Treatment",slug:"controversies-in-pneumothorax-treatment",totalDownloads:481,totalCrossrefCites:0,authors:[{id:"63412",title:"Dr.",name:"Khalid",surname:"Amer",slug:"khalid-amer",fullName:"Khalid Amer"}]}],productType:{id:"1",title:"Edited Volume",chapterContentType:"chapter",authoredCaption:"Edited by"},personalPublishingAssistant:{id:"252211",firstName:"Sara",lastName:"Debeuc",middleName:null,title:"Ms.",imageUrl:"https://mts.intechopen.com/storage/users/252211/images/7239_n.png",email:"sara.d@intechopen.com",biography:"As an Author Service Manager my responsibilities include monitoring and facilitating all publishing activities for authors and editors. 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Panos"}],productType:{id:"1",chapterContentType:"chapter",authoredCaption:"Edited by"}}]},chapter:{item:{type:"chapter",id:"50943",title:"Langmuir‐Blodgett Methodology: A Versatile Technique to Build 2D Material Films",doi:"10.5772/63495",slug:"langmuir-blodgett-methodology-a-versatile-technique-to-build-2d-material-films",body:'\n
Dimensionality is one of the most fundamental material parameters because it defines the atomic structure of materials and determines its main properties. Thus, one chemical element or compound can exhibit different properties in different dimensions. Some interesting examples about size effect are surface plasmon resonance in metal nanoparticles, quantum confinement in semiconductor particles, and superparamagnetism in magnetic nanomaterials.
\nIf one dimension is restricted, we have two‐dimensional (2D) or layered shape material. An interesting example of this kind of materials is graphene. This material is a monolayer of carbon atoms tightly packed into 2D honeycomb lattice that has attracted worldwide attention since it was discovered in 2004 [1, 2]. This new material has emerged with a promising future due to its amazing properties such as transparency, high‐charge mobility, thermal conductivity, and mechanical resistance. Due to these unique properties, graphene has been proposed as a good candidate for manufacturing transparent‐conducting electrodes, transistors, hydrogen‐storage devices, and gas sensors [3, 4]. The growing interest on graphene has highlighted the importance of another 2D material in technological applications such as transition metal chalcogenides (TMC) and layered ionic solids.
\nSeveral 2D materials can be obtained by exfoliation of a layered bulk crystal. However, this procedure is often difficult because the van der Waals interactions between adjacent layers must be overcome. Mechanical exfoliation provides good results, but mostly applied for fundamental research because it is arduous and expensive to produce the material at industrial scale by this way. Other methodologies such as solvent‐assisted ultrasound exfoliation [5] or chemical synthesis [6] allow obtaining large amounts of materials at low cost, although they present several disadvantages against mechanical exfoliation. One important disadvantage is related with the deposition of materials onto solids. Hence, for several technological applications, it is necessary to support 2D materials onto solid substrates [7, 8] and since the properties of 2D materials deposited onto solids are strongly affected by the film morphology, a deposition methodology becomes necessary, which allows a great control of the material density and packing. Several techniques such as drop casting [9] or spin coating [10] have been used to integrate these materials onto novel devices; however, they often lead to nonuniform films or films with aggregated materials due to uncontrolled capillary flow and dewetting processes during solvent evaporation. These aggregates decrease the specific properties of each material [2, 11]. An illustrative example of the aggregation produced by solvent evaporation can be seen in Figure 1. The figure shows a field emission scanning electron microscopy (FESEM) image of a graphene oxide (GO) film deposited onto silicon by drop casting.
\nFESEM image of graphene oxide deposited onto silicon by drop casting.
An alternative technique is the Langmuir‐Blodgett (LB) methodology. This method consists on the transfer process of a water‐insoluble material from the air‐water interface onto a solid substrate by vertical dipping of the solid in the Langmuir monolayer [12]. This technique allows continuous variation of material density, packing, and arrangement by compressing or expanding the film using barriers. Consequently, it offers the possibility of preparing films with the control of interparticle distance necessary to exploit the 2D materials in technological applications. Despite this methodology being successfully used for transferring water‐insoluble molecules [12–14] and nanoparticles [15], it has been less employed to transfer graphene derivatives [16–19] or TMC materials onto solid substrates.
\nThis chapter reviews some strategies to build 2D material films by means of the LB methodology. The content is organized into three main sections. The first one introduces the LB methodology. The second one summarizes the production of thin films of graphene oxide derivatives by using this methodology [17–20]. The last section describes some representative results concerning thin films of Quantum Dots (QDs) of transition metal chalcogenides [21–26] and silver nanowires (AgNWs) [27]. All sections are focused on the possibility of tuning the morphology of the 2D material by modifying the surface composition of the Langmuir monolayer and the deposition methodology.
\nThe LB methodology consists on the transfer process from the air‐water interface onto a solid substrate of a monomolecular layer of amphiphilic material adsorbed at the air‐water interface [12]. The amphiphilic material is dissolved in a volatile solvent and dropped on the air‐water interface. For optimal results, the solvent should have a positive spreading coefficient and be insoluble in the subphase [28]. After the spreading, the solvent evaporates and the material forms a monolayer. When the monolayer reaches the thermodynamic equilibrium, it is symmetrically compressed by using two barriers. The sequential isothermal compression changes the structure of the monomolecular film, which passes through a series of 2D states, referred to as gas, expanded and compressed liquids, and solid state. Consequently, knowing the 2D phase diagram of the film, it is possible to control its structure and associated physical and chemical properties. To transfer the film onto the solid, a flat substrate is immersed into the aqueous subphase and then extracted in a controlled way with the film adsorbed onto it, see Figure 2a. The transfer process can be repeated many times to obtain multilayers [12, 29] of different thickness and composition. During the transfer process, the surface pressure is maintained constant by barrier compression in order to compensate the loss of molecules transferred onto the solid. One variant of this methodology is the horizontal deposition technique, referred as Langmuir‐Schaefer (LS) method. In the latter, the solid substrate is placed parallel to the air‐water interface and deposition is done by contacting the substrate horizontally with the floating monolayer, see Figure 2b.
\nLangmuir‐Blodgett (a) and Langmuir‐Schaefer (b) transfer processes.
When the monolayer is transferred, its structure is often modified; therefore, to construct high‐quality films, a careful control of experimental parameters, such as stability and homogeneity of the monolayer, subphase properties (composition, pH, presence of electrolytes, and temperature), substrate nature, speed of immersion/emersion of substrate, surface pressure during the deposition process, and the number of transferred monolayers, is required [12, 29].
\nAs commented previously, due to its unique properties graphene has been suggested to be used in a great number of technological applications. Nevertheless, each application requires a different set of properties. Thus, graphene synthesized by chemical vapor deposition (CVD) or micromechanical exfoliation renders high‐quality sheets suitable for electronic applications; however, these sheets cannot be used for the fabrication of composites or water‐soluble materials, because they do not contain functionalized groups. In these situations, graphene oxides [30] are preferred because they contain reactive oxygen functional groups that can attach small molecules, polymers, or nanoparticles to the graphitic surfaces for potential use in polymer composites [31], gas sensors [32], or photovoltaic cells [33, 34].
\nAnother important issue of the use of graphene in technological applications is related to its implementation in devices. In the particular case of graphene oxide derivatives, conventional deposition techniques such as drop casting [9] or spin coating [10] not only induce aggregation of flakes, as can be seen in Figure 1, but also force the sheets to fold and wrinkle, losing its excellent properties [2]. Therefore, to overcome these limitations other deposition techniques such as LB have been recently proposed [16, 35].
\nGraphene oxide can be considered as an amphiphilic material [36] because it is constituted by two different domains. The hydrophobic one corresponds to π‐conjugated sp2 carbon while the hydrophilic domain is constituted by O‐groups attached at the basal plane [37]. The existence of two regions allows obtaining stable water‐insoluble monolayers of the material. Accordingly, several properties such as rheological properties, morphology, and stability of GO monolayers have been recently reported [11, 16, 18, 19]. Several works seem to indicate a great influence of the chemical composition on the film properties [19, 38].
\nConcerning the chemical synthesis, graphene oxide is usually obtained by graphite oxidation [17, 19, 39] or carbon nanofibers [18–20, 40] by means of the Staudenmaier [41] or the Hummers [42] reactions. Then, graphene oxides are often reduced by chemical agents [17, 43, 44] or thermal annealing [45, 46] to restore the graphene structure. However, both the GO reduced by chemical agents, referred as reduced graphene oxide (RGO), and the thermally reduced retain some O‐groups attached to the basal plane of GO. These O‐groups decrease the amazing properties of graphene such as transparency and high electric conductivity.
\nDespite the interest raised by GO, the knowledge of its chemical structure remains still a challenge. The best‐known graphene oxide structure consists of two different carbon domains constituted by Csp2 corresponding to aromatic groups and Csp3 of alcohol and epoxy groups attached at the basal plane. The carboxylic acid groups are located at the edge of the sheets [37]. However, the main origin of the controversy is the percentage of each group into the flakes. Several are the causes of discrepancies, although the variability of the starting material and the oxidation process seem to be the most important ones [47]. On the other hand, the chemical structure of graphene oxide was recently revisited because it has been proved that the oxidation of carbon‐based materials originates highly oxidized fragment, named as oxidative debris (OD) [47–49]. The oxidized fragments remain strongly adsorbed onto the graphitic sheets due to π‐π staking interactions but can be removed by alkaline washing of graphene oxide. The purified GO contains lower O/C ratio than the non‐purified one, and consequently its chemical structure and solubility properties are quite different [47, 49].
\nIt is necessary to consider that in nanocomposites built with GO, the second component, polymers, nanoparticles, or small molecules, often interacts with the O‐groups of graphene oxide; therefore, to improve the quality of the composite, it is crucial to have knowledge of the chemical structure of graphene oxide to control interactions between components which have a great influence on the properties of nanocomposites. However, there is no systematic study related to the effect of the oxidation procedure, nature of the starting material, and purification process on the chemical structure and properties of graphene oxides. Recently, we have started the systematic study of the effect of the starting material, reduction protocol, and purification process on the chemical structure of graphene oxides and on the film morphology. With this objective in mind, we have synthesized graphene oxides using graphite, and GANF® nanofibers from the Grupo Antolín Ingenieria (Burgos, Spain) as starting materials. The oxidation procedure was Hummer\'s reaction modified to obtain more oxidized samples [17, 18–20]. As reducing agents, we used hydrazine, vitamin C, and sodium borohydride. The purification process consisted of alkaline washing of graphene oxide and is previously reported [48, 49].
\nTo quantify the oxidation degree of different materials, X‐ray photoelectron spectroscopy (XPS) was employed. In all samples, the C1s core‐level spectrum is an asymmetric band that can be fitted to three components centered at 284.8, 286.4, and 287.9 eV. These peaks are assigned to C-C bonds of the aromatic network, C-O bonds of alcohol or epoxide groups, and COOH groups, respectively [50]. From the area of these peaks, the percentage of the different groups in each sample was calculated. Results obtained for different kinds of graphene oxides are collected in Figure 3. Data shown in Figure 3 were taken from references [17–19].
\nChemical composition of graphene oxides determined by XPS measurements. Data were taken from references [17–19]. Solid symbols correspond to graphene oxides obtained from graphite and open symbols from GANF® nanofibers. Stars are results of surfactant‐functionalized graphene oxides.
Results in Figure 3 clearly show differences between the chemical composition of graphene oxides synthesized by the oxidation of graphite (GO) and nanofibers (NGO). Thus, the percentage of Csp2 is slightly higher for NGO than for GO, while the percentage of C-O groups at the basal plane is higher for GO than for NGO and the percentage of COOH groups attached to NGO is twice that of GO. This behavior was attributed to the different size of nanoplatelets [18, 19]. In the case of NGO, dynamic light‐scattering measurements (DLS) and the statistical analysis of FESEM images demonstrated that nanoplatelets of NGO are smaller than the GO ones; therefore, since the carboxylic groups are mainly localized at the edge of sheets the smallest sheets contain the highest proportions of COO- groups [19]. As far as the influence of the purification procedure on the chemical composition, our results indicated that the percentage of Csp2 increases after the alkaline washing. Moreover, the purification process drives to graphene oxides of similar chemical composition although the chemical structure of non‐purified graphene oxides is quite different.
\nAnother interesting result is that the percentage of Csp2 of reduced graphene oxide is almost independent on the reducing agent, and the averaged value of 65 ± 2 is lower than the value found for purified graphene oxide, 72 ± 4. This fact was previously reported for graphene oxide reduced by hydrazine [47] and was interpreted as follows: due to the basic nature of hydrazine, it cleans oxidative debris and simultaneously reduces the oxygen groups of graphene oxide; however, nitrogen atoms remain attached to RGO sheets decreasing the percentage of Csp2. This C-N bond identified by XPS as a peak centered at 400 eV is responsible for the increase of Csp3 percentage. The balance of these processes leads to RGO sheets of intermediate composition between purified PGO and GO [17]. Similar situations were observed for graphene oxides reduced by vitamin C and borohydride, respectively. In these cases, oxygen and boron atoms of the oxidized product of vitamin C and borohydride remain attached to the network decreasing the aromatic degree of graphene derivatives. According to our results, we postulate that alkaline washing must be the preferred procedure to increase the Csp2 percentage on graphene oxide nanoplatelets.
\nGraphene oxide nanoplatelets are insulators and to increase the electric conductivity chemical reduction has been postulated. However, RGO films prepared by conventional deposition methodologies present low electrical conductivity values. This is probably due to the platelet aggregation induced by dewetting processes. We have explored the Langmuir‐Blodgett methodology to obtain non‐aggregated and ordered reduced graphene oxide films. To prepare the LB film, it is necessary to select the proper surface state, which will be transferred. To identify the surface state of materials at the interface, the compressional modulus, ε, has been widely used. The parameter can be calculated from the surface‐pressure isotherm using Eq. (1):\n
In Eq. (1), A represents the surface area and π the surface pressure value. We have recorded the surface pressure isotherms of each material and a representative example is plotted in Figure 4a. The compressional elastic modulus value is plotted against the surface pressure in Figure 4b.
\nThe isotherm morphology is similar to that of surfactant molecules and can be interpreted as follows: monolayers of surface pressure close to zero correspond to low values of compressional modulus and were assigned to surface states in which the nanoplatelets are isolated in a two‐dimensional gas state. When the surface area decreases, the nanoplatelets are pushed closer to each other, resulting in small domains in which ε grows until it reaches a maximum value. This two‐dimensional region is commonly assigned to the liquid‐expanded (LE) state and corresponds to close‐packed sheets. Beyond the compressional elastic modulus maximum, the nanoplatelets form wrinkles, overlaps, and three‐dimensional (3D) structures [16].
\n(a) Surface pressure and (b) compressional elastic modulus isotherms of graphene oxide reduced by borohydride at 293 K.
In a previous work, the LE state of the GO monolayer [38] has been modeled by Volmer\'s model adapted to nanoparticles [51]. We have used this model to interpret the isotherms of different nanoplatelets of GO at the LE state. Our results demonstrated strong interactions between carboxylic acids at the edge of sheets through hydrogen bonds [18, 19].
\nBecause we are interested to build GO films of closely packed and nonoverlapped nanoplatelets, we transferred graphene oxide monolayers at the LE state by the LB methodology [18, 19]. Representative atomic force microscopy (AFM) and FESEM images of these films are collected in Figure 5.
\nRepresentative images of different graphene oxides films: (a) SEM image of graphene oxide obtained by oxidation of graphite; (b) TEM image of graphene oxide reduced by vitamin C. The inset is a magnification to show the morphology of RGOv nanoplatelets; (c) graphene oxide functionalized with DDPS and reduced by hydrazine. The inset shows a magnification of the AFM image; (d) graphene oxide functionalized with DDPS and reduced by vitamin C. The inset is a TEM image to show details of the film morphology. Reduced graphene oxides were obtained using graphite as starting materials. The surface pressure of the Langmuir monolayer precursor of the LB film was 1 mN m-1.
As can be seen in Figure 5, the solid coverage is higher for GO, Figure 5a, than for reduced graphene oxides, Figure 5b. Low coverage was also reported for purified graphene oxides [18, 19] and was attributed to the low percentage of O‐groups attached to sheets of purified oxides [19]. A high percentage of O‐groups favor the contact between silanol groups of the silicon wafer and sheets increasing the adhesion of nanoplatelets to silicon. Since the chemical composition of reduced and purified graphene oxides is almost the same, the low percentage of O‐groups at reduced samples can be responsible for the low coverage observed for RGO films.
\nIn an attempt to improve the solid coverage, the reduced graphene oxides were functionalized with the zwitterionic surfactant N dodecyl‐N,N‐dimethyl‐3‐ammonio‐1‐propanesulfonate (DDPS). We have proved that the surfactant remains adsorbed onto graphene oxide platelets playing two important roles: as surface active molecule, it favours attractive interactions between the silicon and the reduced graphene oxide and because it is attached at sheets minimizing the restacking of flakes. It is interesting to note that the surfactant is attached to sheets in a non‐covalent way, and consequently the chemical structure of graphene oxide is not significantly altered [52].
\nThe AFM images of functionalized reduced graphene oxide films, Figure 5c and d, show that the functionalization with the DDPS surfactant increases the solid coverage; however, it is lower than that for graphene oxide, Figure 5a. The AFM images of RGOhS, Figure 5c, also show the formation of the chained sheets suggesting lateral attractive interactions between flakes. These attractive interactions can be likely induced by the surfactant molecules attached to the sheets [17].
\nWe have great interest to study the effect of GO chemical composition on the electrical conductivity of GO films. However, in the case of reduced graphene oxide the electrical conductivity value is too small to detect significant differences; therefore, we employed an alternative method widely used by other authors. The method consists of measuring the conductivity of paper‐like graphene oxide films [53]. To analyze the electrical conductivity dependence with the chemical composition, we have plotted the electrical conductivity against the Csp2 and C-O group percentages shown in Figure 6a and b, respectively.
\nVariation of the electrical conductivity values of paper‐like graphene oxide films with the Csp2 (a) and (b) C-O group percentages, respectively. Data were taken from Reference [16].
Results in Figure 6a show that the electrical conductivity increases as the Csp2 percentage. Moreover, the highest conductivity value is obtained for graphene oxide functionalized with the zwitterionic surfactant. In addition, samples with the lowest percentage of C-O and COOH groups, see Figure 3, correspond with reduced graphene oxides functionalized with the surfactant DDPS. All these facts suggest that the surfactant molecules can eliminate high amount of O‐groups of samples increasing the electrical conductivity of flakes as can be seen in Figure 6b.
\nOn summarizing, the LB technique can be presented as a good methodology of building graphene oxide films because it renders high‐coverage and ordered films. On the other hand, the conductivity of our surfactant‐functionalized RGO samples is higher than the values found in the literature for paper‐like films of reduced graphene oxide [5, 54] functionalized with ionic surfactants, although more efforts must be done to improve the solid coverage and to increase the electrical conductivity values of graphene oxide films.
\nNanoparticles of CdSe Quantum Dots are semiconductors which show size dependence in their optoelectronic properties with attractive applications in the fabrication of solar cells or light‐emitting diodes (LEDs) due to their band‐gap tunability.
\nThe most important optical advantages are a broad and continuous absorbance spectrum and a narrow emission spectrum whose maximum position and dynamic emission properties depend on its QD size. However, optoelectronic device applications based on nanoparticles require QDs assembly in controllable architecture to avoid the deterioration of the quantum film efficiency. Therefore, the thickness and uniformity of the assembled QD films are crucial factors in the emission properties of films [7, 55–58].
\nIn the particular case of CdSe QDs, the hydrophobic nanoparticles present the highest quantum efficiency. However, when these nanoparticles are transferred from the air‐water interface onto substrates such as glass, silicon, or mica without treatment to become the solid surface, hydrophobic, low coverage and nanoparticle agglomeration have been observed [59, 60]. These undesirable results decrease the quantum yield of nanoparticle films. To solve this problem, some approaches have been proposed. One of the most widely used strategies consists of mixing nanoparticles with surfactants or polymers and then transferring the mixture from the air‐water interface onto the solid substrate [61–63]. This approach seeks to control the assembly of hydrophobic nanoparticles at the air‐water interface. With this purpose, we have proposed three amphiphilic molecules of distinct nature, the copolymers poly(octadecene‐co‐maleic anhydride), PMAO, and poly(styrene‐co‐maleic anhydride) partial 2‐butoxyethyl ester cumene terminated, PS‐MA‐BEE, and the Gemini surfactant ethyl‐bis(dimethyl octadecylammonium bromide), 18‐2‐18. All these molecules present surface activity and can anchor to substrates such as mica, glass, or silicon, through their hydrophilic moieties [13, 14] favoring the QDs’ adhesion across its hydrophobic part. We have chosen the polymer PMAO because it interacts effectively with hydrophobic nanoparticles leading to excellent stability by avoiding 3D aggregation [64, 65]. In the case of the polymer PS‐MA‐BEE, it was chosen because it is a good component to organize hybrid nanomaterials used in submicrometric electronic devices [66]. This is due to its mechanical rigidity and good adhesion on solids [67]. Finally, the Gemini surfactant was chosen since it has been proposed in combination with DNA for biotechnological applications [68, 69].
\nOur results demonstrated that the QD aggregation is avoided by the addition of these polymer and surfactant molecules. Attractive interactions between the chains of these molecules and the hydrophobic moieties of the QD stabilizer, trioctylphosphine oxide (TOPO), favor the adsorption of QDs on the matrices, while the hydrophilic groups of polymer or surfactant molecules increase the QDs’ adhesion in solid substrates, avoiding the nanoparticle agglomeration.
\nWe also found two different film features depending on the film composition. To illustrate this behavior, Figure 7 collects some AFM, transmission electron microscopy (TEM), and scanning electron microscopy (SEM) images of QD films prepared with different matrix compositions.
\n(a) AFM image of a Gemini/QD film at the surfactant mole fraction of 0.98, (b) SEM image of a PMAO/QD LB film onto mica at the polymer mole fraction of 0.50. The Langmuir monolayers were transferred at the surface pressure of 30 mN m-1, (c) TEM image of mixed PS‐MA‐BEE/QD LB film of polymer mole fraction of 0.5 and deposited at the surface pressure of 14 mN m-1, and (d) AFM image of a mixed PS‐MA‐BEE/QD film of polymer mole fraction of 0.96 and deposited at the surface pressure of 30 mN m-1. The inset corresponds to the TEM image of the film.
Images in Figure 7 show two different morphologies, hexagonal networks and domains of different shapes, depending on the film composition. Thus, mixed films of QDs and PS‐MA‐BEE of high polymer mole fraction, XP ≥ 0.95, and deposited at 30 mN m-1 [22] and PMAO/QDs films are constituted by hexagonal networks [21], Figure 7b and d. It is interesting to note that the height of rims around the holes was 4 nm, which is compatible with the diameter of the nanoparticles dissolved in chloroform (3.4 nm). This result indicates that QDs are mainly confined in rims and do not form 3D aggregates. On the other hand, all the Gemini/QDs films and PS‐MA‐BEE/QDs films of polymer mole fraction below 0.95 deposited at low‐surface pressure (14 mN m-1) are constituted by domains of different morphologies, Figure 7a and c. The domain height determined by AFM measurements (∼3 nm) is consistent with the diameter of QDs dissolved in chloroform. This fact indicates that there is no 3D aggregation in these films.
\nDifferences between film morphologies were interpreted according to dewetting mechanisms [21, 22]. The two dewetting mechanisms considered in these cases are known as nucleation, growth, and coalescence of holes [70] and spinodal [71]. In the former, the gravity contribution predominates and the dewetting process starts with the nucleation of holes at film‐defect sites followed by the material displacement away from the nucleus. The material is accumulated in the rims of holes delimiting a mosaic [70]. Conversely, in the spinodal dewetting mechanism, the capillary waves break the film into nanostructures when the amplitude of the capillary waves exceeds the thickness of the film. Taking into account that the molecular weight of the polymer PMAO is around 50 times higher than the surfactant one, it becomes clear that the gravitational effect prevailed over the capillary waves even in films with small amount of the polymer PMAO. Therefore, the PMAO/QDs film morphology is driven by the mechanism of nucleation, growth, and coalescence of holes, while spinodal dewetting mechanism prevails in Gemini/QD films [21]. In the case of PS‐MA‐BEE/QD films, the interpretation of the behavior observed is not so evident and it is necessary to analyze the balance between the driving forces involved in the surface arrangement: gravitational and capillary forces. Thus, the elasticity values go through a minimum for PS‐MA‐BEE/QDs monolayers at the surface pressure of 30 mN m-1 and for polymer mole ratio above 0.95, while it reaches maximum values for monolayers at the surface pressure value of 14 mN m‐-1 and XP < 0.95 [22]. Taking into account that the damping coefficient passes through a maximum at low elasticity values and decreases when the elasticity modulus increases [22], it is easy to understand that in PS‐MA‐BEE/QD films of low elasticity values (π = 30 mN m-1 and XP ≥ 0.95), the capillary waves are quickly damped and the film breaks in domains separated by holes due to gravitational effects. Conversely, the capillary waves for monolayers with the highest elasticity values (π = 14 mN m-1 and XP < 0.95) do not damp so quickly and they drive the dewetting mechanism. In these situations, the spinodal dewetting mechanism predominates against the growth of holes process leading to QD domains of different shapes [22].
\nAnother interesting example is the preparation of silver nanowire films for manufacturing modern devices such as photovoltaic cells, touch panels, and light‐emitting diodes. Although the development of new materials is mainly by the requirements of each application [72], high transparency and electrical conductivity always constitute required requisites.
\nIndium tin oxide (ITO) currently dominates the field of transparent conductive electrodes as a result of its excellent optoelectronic properties [73]; however, it suffers important limitations due to the scarcity of indium, brittleness of its electrodes, and high manufacturing cost. Several materials such as carbon nanotubes [74, 75], graphene films [76, 77], conducting polymers, and metal nanowires [72, 78] are being analyzed to replace ITO. However, the properties of these materials, in terms of electrical resistance and transparency, are still inferior to ITO [78]. Among all, the silver nanowires arouse great interest due to the high conductivity of silver (6.3 × 107 S m-1) [79]. Since the nanowires are usually synthesized in solution, an important issue is the control of the transfer process from solutions onto the substrate. This is because to achieve low electrical resistance and high transparency, it is necessary to optimize the morphology, the placement of nanowires, and the junction resistance between them in the network. As commented previously, spin‐coating and drop‐casting methodologies present several disadvantages since water evaporation leaves discontinuous films with typical coffee rings that significantly decrease the quality of AgNW films [80, 81]. To overcome these limitations, we have reported a strategy based on the Langmuir‐Schaefer methodology to transfer hydrophobic AgNW from the air‐water interface onto Lexan polycarbonate substrate in an ordered and controlled way [27].
\nThe first step for building LB films is to obtain stable monolayers of hydrophobic materials. Therefore, it is necessary to synthesize hydrophobic nanowires, since the commercial ones are water soluble since they use polyvinyl pyrrolidone molecules as capping agents. To synthesize hydrophobic AgNW, polyvinyl pyrrolidone was replaced by octyl thiol molecules. The surface modification is achieved through the surface ligand exchange procedure reported by Tao [82]. After the synthesis of AgNW, they were deposited at the air‐water interface and different surface states were transferred onto the solid substrate by the LS methodology. The surface states of nanowire monolayers are characterized by the surface compressional modulus, ε, calculated from the surface pressure isotherm and Eq. (1), and ε‐values are plotted against the surface concentration, Γ, in Figure 8a. As can be seen in Figure 8a, when the surface concentration is small, the elasticity modulus value is close to zero. In this region, named as low‐surface density state (LD), nanowires are randomly orientated. When the surface density is further increased and ε reaches a value of 10 mN m-1, the monolayer is highly packed; we referred to this state as the high‐surface density state (HD) [83].
\n(a) Elasticity isotherm of silver nanowires capped with octyl thiol at 20°C, (b) FESEM image of a bilayer of AgNW of 645 mg m-2. Arrows indicate the orientation of the first (red) and second (blue) layers, and (c) variation of sheet resistance and transmittance with the nanowire surface concentration of LS films.
We have transferred AgNW Langmuir monolayers at LD and HD states by the Langmuir‐Schaefer methodology. With the purpose of achieving a network of nanowires, a second layer in which the nanowires are oriented perpendicular to the first layer was deposited. In the first and second layers, the surface density of the transferred Langmuir monolayer was the same [27]. The surface density is controlled by the surface pressure value. Figure 8b shows a representative FE‐SEM image of a nanowire film obtained by this methodology.
\nThe sheet resistance, Rs, measured in Ω sq-1 and the transmittance measured at 550 nm are plotted against the surface concentration in Figure 8c. Data in Figure 8c show that the monolayers at the LD state present high Rs values which decrease when the surface concentration increases, while the transmittance value is almost independent on surface concentration and remains constant at 88%. The behavior is opposite for films built from Langmuir monolayers at the HD state. In this case, the sheet resistance is maintained at 8 Ω sq-1 while the transmittance value changes from 65 to 89% when the surface concentration was modified between 345 to 770 mg m-2. According to the resistance and transparency values, our AgNW films can be employed as substitutes for ITO as components of devices such as touch screens, electromagnetic shielding, and defrosted windows [27]. Moreover, our results proved that the Langmuir‐Schaefer methodology is a versatile technique, which allows modifying the transmittance keeping the sheet resistance or tuning the sheet resistance, maintaining the transparency of films constant by properly selecting the surface state and the nanowire mass transferred onto the solid substrate.
\nResults analyzed in this chapter allow us to discuss the ability of the Langmuir‐Blodgett and Langmuir‐Schaefer methodologies to build thin films of 2D materials such as graphene oxides, transition metal chalcogenide nanoparticles, CdSe Quantum Dots, and silver nanowires. We discuss the advantages of these methods against the most conventional ones such as drop and spin coating for built‐in 2D material films with applications in the fabrication of solar cells, LEDs, sensors, and transparent electrodes.
\nWe also review some strategies for improving the solid coverage, avoiding the nanoparticle aggregation, and modulating the film morphology. All these issues are crucial for increasing the quality of films and to modulate its properties according to the properties required for each application.
\nResults analyzed in this chapter indicate that the Langmuir‐Blodgett and Langmuir‐Schaefer methodologies combined with self‐assembled materials can be proposed as a non‐template reproducible technique for patterning at the nanoscale. However, most efforts have to be done for achieving more homogeneous films, higher coverage, and a greater control of the material arrangements to build good‐quality films to be used in technological applications.
\nThe authors thank financial support from the European Regional Development Fund, ERDF, Ministerio de Educación y Ciencia (MAT 2010‐19727), and Ministerio de Economía y Competitividad (IPT‐2012‐0429‐420000). TA and BMG wish to thank the European Social Fund and Consejería de Educación de la Junta de Castilla y León for their FPI grants. We also thank Ultra‐Intense Lasers Pulsed Center of Salamanca (CLPU) for the AFM measurements, to Microscopy Service (Universidad de Salamanca) for the TEM measurements, and Sala Blanca de Nanotecnología (USAL) for FE‐SEM facility. We thank Dr. García Fierro (Instituto de Catálisis y Petroleoquímica, Madrid) for XPS measurements.
\nThe 354 alloy belongs to the Al-Si-Cu-Mg system similar to B319 alloy that is widely used for automotive engine blocks [1]. The high silicon content in the 354 alloy improves the alloy castability whereas the presence of Cu and Mg noticeably enhances the yield strength (YS) and the ultimate tensile strength (UTS) of the 354 alloy due to the formation of intermetallic phases, mainly Al2Cu or eutectic Al + Al2Cu, and Mg2Si precipitates [2, 3]. However, segregation behavior of Cu may lead to incipient melting during solution treatment which will apparently reduce the alloy strength [4]. Addition of Mg has a strong affinity to react with Sr, leading to the formation of a complex Mg2SrAl4Si3 intermetallic phase, and hence reducing the effectiveness of Sr as a Si modifying agent [5]. In the absence of Cu, high Fe and Mg contents lead to the formation of π-FeMg3Si6Al8 phase which is difficult to dissolve during the solution treatment process [6, 7]. In the quaternary Al-Si-Cu-Mg alloy system, Q-phase (Al4Mg8Cu2Si6) can coexist with the Al2Cu, Mg2Si, and Si phases depending on the levels of Cu, Mg, and Si [8, 9, 10, 11]. The different factors that may influence the mechanical behavior of cast aluminum alloys are schematically represented in Figure 1 [12].
\nSchematic representation of factors affecting alloy performance [12].
Zirconium may be added to Al alloys in order to refine the grain structure due to the presence of fine coherent dispersoids (mainly Al3Zr) which obstruct dislocation motion and in turn, enhance the elevated temperature mechanical properties of aluminum alloys [13]. In order to increase the volume fraction of Al3Zr precipitates and based on the phase diagram of Al-Zr, the concentration of Zr in the alloys investigated in this study was kept at around 0.3 wt.% [14].
\nThe main purpose of solution heat treatment is to obtain a supersaturated solid solution at high temperatures (below the eutectic temperature). As a result, a homogeneous supersaturated solid solution (SSSS) will form through dissolving the precipitated phases during the solidification process, such as β-Mg2Si, θ-Al2Cu, Q-Al5Cu2Mg8Si6, π-Al9FeMg3Si5 and β-Al5FeSi phases. The β-Mg2Si and θ-Al2Cu phases can be easily dissolved when the optimum solution heat treatment temperature and time are employed. The solution treatment temperature is determined according to the alloy composition and solid solubility limit; however, it must be lower than the melting point of the phases that exist in the as-cast structure to avoid incipient melting of these phases [15, 16].
\nFollowing the empirically developed concept of quality index proposed by Drouzy et al. [17, 18] Cáceres proposed a mathematical model emphasizing the significance of the quality index as follows [17, 19, 20]:
\nwhere quality index Q can be calculated using the relative quality index (q), strain-hardening exponent (n), and the strength coefficient (K).
\nThe present study was undertaken to explore the effect of Zr addition and aging conditions of the as cast tensile bars on:
Characterizing the microstructural features of the investigated alloys,
Exploring the tensile properties at both ambient and elevated temperatures, and
Correlating the tensile properties to the microstructural features to establish the strengthening or softening mechanisms responsible for the observed properties.
It should be noted here that the term “temperature” applies to aging temperatures as well as testing temperature.
\nAlloy 354 modified with 200 ppm of strontium (using Al-10% Sr master alloy) and grain refined using 0.20 wt.%Ti (Al-5%Ti-1%B) was used as the base alloy (alloy A). To this alloy, 0.3%Zr in the form of Al-25wt.%Zr master alloy was added (alloy B). The chemical compositions of both alloys are listed in Table 1. Figure 2 shows the microstructure of the as-received base alloy ingots. Melting and casting procedures were carried out as described elsewhere.
\nMicrostructure (200×) of the base alloy 354 used in this work.
To prepare test bars for the tensile tests, three samples for chemical analysis were also taken at the time of the casting; this was done at the beginning, in the middle, and at the end of the casting process to ascertain the exact chemical composition of each alloy. The experimental work was divided into two stages: Stage I in which the 354 alloy (alloy A) was used, and Stage II where the 354 alloy with 0.3%Zr (alloy B) was used. In Stage I, the melt temperature was kept around 750°C, whereas in Stage II, the melt temperature was superheated to 800°C, to ensure the complete decomposition of the Al-25%Zr master alloy used.
\nTensile bars were solution heat treated at 495°C for 8 h, followed by quenching in warm water at 60°C, after which artificial aging was applied according to the plan listed in Table 2. After aging, the test bars were allowed to cool naturally at room temperature (25°C). All of the samples, whether as-cast, solution heat-treated, or aged, were tested to the point of fracture using an MTS servo-hydraulic mechanical testing machine at a strain rate of 4 × 10−4 s−1.
\n\nAlloy code | \nElement (wt.%) | \n||||||||
---|---|---|---|---|---|---|---|---|---|
Si | \nFe | \nCu | \nMn | \nMg | \nZr | \nTi | \nSr | \nAl | \n|
A | \n9.1 | \n0.12 | \n1.8 | \n0.0085 | \n0.6 | \n— | \n0.18 | \n0.02 | \n87.6 | \n
B | \n9.1 | \n0.12 | \n1.8 | \n0.0085 | \n0.6 | \n0.3 | \n0.18 | \n0.02 | \n87.6 | \n
Chemical composition of the 354 alloys used in this study.
Temperature (°C) | \nAging time (h) and aging condition codes | \n||||||||||||
---|---|---|---|---|---|---|---|---|---|---|---|---|---|
2 | \n4 | \n6 | \n8 | \n10 | \n12 | \n16 | \n20 | \n24 | \n36 | \n48 | \n72 | \n100 | \n|
155 | \n1 | \n2 | \n3 | \n4 | \n5 | \n6 | \n7 | \n8 | \n9 | \n10 | \n11 | \n12 | \n13 | \n
170 | \n14 | \n15 | \n16 | \n17 | \n18 | \n19 | \n20 | \n21 | \n22 | \n23 | \n24 | \n25 | \n26 | \n
190 | \n27 | \n28 | \n29 | \n30 | \n31 | \n32 | \n33 | \n34 | \n35 | \n36 | \n37 | \n38 | \n39 | \n
240 | \n40 | \n41 | \n42 | \n43 | \n44 | \n45 | \n46 | \n47 | \n48 | \n49 | \n50 | \n51 | \n52 | \n
300 | \n53 | \n54 | \n55 | \n56 | \n57 | \n58 | \n59 | \n60 | \n61 | \n62 | \n63 | \n64 | \n65 | \n
350 | \n66 | \n67 | \n68 | \n69 | \n70 | \n71 | \n72 | \n73 | \n74 | \n75 | \n76 | \n77 | \n78 | \n
Artificial aging conditions used for room temperature tension tests.
The yield strength (YS) was calculated according to the standard 0.2% offset strain, and the fracture elongation was calculated as the percent elongation (%El) over 50 mm gauge length, as recorded by the extensometer. The ultimate tensile strength (UTS) was also obtained from the data acquisition system of the MTS machine. The average %El, YS, or UTS values obtained from the five samples tested per condition were considered to be the values representing that specific condition. An extensometer, or strain gage was used in the tests to measure the extent of deformation in the samples.
\nSamples for metallography were sectioned from the tensile-tested bars of all the alloys studied, about 10 mm below the fracture surface. The percentage porosity and eutectic Si-particle characteristics were measured and quantified using an optical microscope linked to a Clemex image analysis system. The microstructures of the polished sample surfaces were examined using an Olympus PMG3 optical microscope. Phase identification was carried out using electron probe microanalysis (EPMA) in conjunction with wavelength dispersive spectroscopic (WDS) analysis, using a JEOL*JXA-889001WD/ED combined microanalyzer operating at 20 kV and 30 nA, where the electron beam size was ~2 μm.
\nMapping of certain specific areas of the polished sample surfaces was also carried out where required, so as to show the distribution of different elements within the phases. The fracture surfaces of tensile-tested samples were also examined using the same SEM, employing the backscattered electron (BSE) detector and EDS system. The fracture behavior was analyzed using the backscattered electron (BSE) images obtained, and analysis of the EDS spectra of phases observed on the fracture surface. Differential scanning calorimetry (DSC) was used to characterize the sequence of reactions occurring during the heating and/or cooling cycles of an alloy sample during a DSC scan which continuously changes with the increasing or decreasing temperature cycle to produce peaks according to the two expected reactions:
Phase formation → heat emission → exothermic peak
Phase dissolution → heat absorption → endothermic peak
For the high temperature tensile tests, samples from selected conditions were tested to fracture using an Instron Universal mechanical testing machine at a strain rate of 4 × 10−4 s−1. The heating furnace installed on the testing machine is an electrical resistance, forced-air box type, having the dimensions 30 × 43 × 30 cm. The yield strength (YS) was calculated according to the standard 0.2% offset strain, and the fracture elongation was calculated as the percent elongation (%El) over the 25.4 mm gauge length as recorded by the extensometer. The ultimate tensile strength (UTS) was obtained from the data acquisition system of the universal machine. In order to reach and stabilize the intended test temperature during the tests, at the time that the samples were mounted in the tensile machine, the furnace was already pre-set at the required temperature; also, these samples were kept mounted in the furnace of the tensile testing machine for 30 min before the start of every test.
\n\nFigure 3 shows the macrostructure revealing the grain size for alloy A, about 200 μm. A complete modification of the silicon particles in the microstructure of alloy A in the as-cast condition can be seen in Figure 4(a). From Figure 4(a) and (b), solution heat treatment has changed the morphology of the silicon particles from faceted to globular. As a consequence of solution heat treatment, there may also be observed a reduction in the number of silicon particles and a reduction in the density of the silicon phase, due to the diffusion of silicon into the aluminum matrix. The white arrows in Figure 4(a) show the rounded shape of the dendrites with grain refining [21], whereas Figure 4(b) reveals the dissolution of the Al2Cu phase observed in Figure 4(a)—circled.
\nMacrograph showing grain size of the tensile bars in the as-cast condition.
Optical microstructure: (a) before, and (b) after solution heat treatment.
Zhu and Liu [22] proposed a model of the granulation of unmodified eutectic Si composed of three major stages during heat treatment: (i) the mass transport of solute, (ii) a discontinuous phase fragmentation, and lastly (iii) spheroidization. During heat treatment, the silicon atoms in the matrix at the Si particle tips diffuse to locations on the curved surfaces of the particles, leading to the dissolution of eutectic silicon at the tips. This transport of silicon atoms ultimately causes the fragmentation and spheroidization of eutectic silicon which is important from strength point of view compared to Si particles with sharp edges which act as sites for stress concentration.
\nThe values of secondary dendrite arm spacing (SDAS), porosity, modification level, and grain size for both the as-cast (AC) and solution heat-treated (SHT) condition are listed in Tables 3 and 4. As can be seen, SHT resulted in (i) no noticeable change in both the SDAS and grain size, (ii) a significant decrease in the particle density due to coarsening of the eutectic Si particles, and (iii) almost complete solubility of Al2Cu in the aluminum matrix. Since the solutionizing temperature was well below the incipient melting temperature, tensile test bars revealed negligible change in the amount of porosity, i.e., no incipient melting.
\nAlloy code condition | \nSDAS (μm) | \nGrain size (μm) | \nPorosity (%) | \nVolume fraction of intermetallics (%) | \n||
---|---|---|---|---|---|---|
EPMA | \n||||||
Av\n*\n\n | \n\n | Av | \nSD\n**\n\n | \nAv | \nSD | \n|
A-AC | \n19.3 | \n201 | \n0.14 | \n0.06 | \n3.08 | \n0.32 | \n
A-SHT | \n23.1 | \n192 | \n0.12 | \n0.05 | \n1.27 | \n0.10 | \n
SDAS, porosity%, grain size, level of modification, and volume fraction of intermetallics for alloy A.
Average.
Standard deviation.
Alloy code condition | \nArea (%) | \nParticle length (μm) | \nRoundness ratio (%) | \nAspect ratio | \nDensity (particles/mm2) | \n|||
---|---|---|---|---|---|---|---|---|
Av | \nAv | \nSD | \nAv | \nSD | \nAv | \nSD | \n||
A-AC | \n14.58 | \n3.522 | \n3.94 | \n0.4302 | \n0.181 | \n2.033 | \n0.8609 | \n39110 | \n
A-SHT | \n10.868 | \n4.286 | \n3.144 | \n0.554 | \n0.1547 | \n1.641 | \n0.5429 | \n12080 | \n
Silicon particle characterization for alloy A.
\nFigure 5 illustrates the effect of aging treatment on the alloy strength parameters. The main observations inferred from this figure can be summarized as follows:
Solution heat treatment and artificial aging at 190°C for 2 h or at 155°C for 100 h resulted in an increase in the alloy strength by ~64% over its as-cast strength.
Aging at 155 or 170°C for a long period of time offered maximum resistance to softening.
The greatest decrease in tensile strength occurred at 240°C (312 MPa at 2 h to 240 MPa at 100 h). Similarly, a significant decrease in strength took place upon aging at 190°C for a lengthy time (from 382 MPa at 2 h to 314 MPa at 100 h) indicating the end of peak-aging or the commencement of over aging.
The greatest reduction in the alloy UTS and YS levels resulted when the tensile bars were aged at a temperature as high as 350°C even for a short period of 2 h.
In comparison to the ascendant and steady strength curves corresponding to aging temperatures of 155°C and 170°C, fluctuations in the strength curves were observed at aging temperatures of 190°C and over, similar to that reported by Reif [23] where a similar alloy was used and an ascendant monotonic strength curve was observed at an aging temperature of 155°C.
Although the highest ductility values were obtained after 2 h aging at 350°C (~5%), none of the aging conditions reached the higher ductility values exhibited by the solution heat-treated condition (~6.5%). This observation may be considered evidence that the mechanical behavior displayed by alloy A is common to that of the Al-Si-Cu-Mg alloys whose strength is obtained at the expense of ductility [24, 25].
Variation in alloy tensile parameters as a function of aging temperature and time: (a) UTS, (b) YS, and (c) %El.
In order to analyze the alloy quality by means of the Quality Index charts, the as-cast and the solution heat treated conditions plus aging conditions at 155°C, 190°C, and 350°C for aging times in the range of 2–100 h were used. From a previous study [5], K was calculated as 500 MPa.
\nThe plastic strain and the quality index (Q ) both exhibit a great improvement following solution heat treatment. The fact that plastic deformation (q) was about 0.31 in the solution-treated condition means that the alloy reached 31% of its maximum quality index value (Q ). The importance of q is that it shows how much a sample is away from its maximum possible ductility q = 1 and indicates that it would be possible to control the microstructure, for example by reducing the SDAS, or the porosity, or intermetallic level to enhance the alloy ductility and hence, the quality index, Q. When the ductility increases sharply from the as-cast to the solution heat treated condition, such changes can be related to the spheroidization of silicon particles and to the uniformity of the microstructure in the solution heat-treated condition, as shown in Figure 6(a).
\n\nQ-charts following: (a) SHT, (b) aging at 155°C, (c) aging at 190°C, and (d) aging at 350°C. Legends in (b) apply for other charts. The curved lines indicate the passage from 2 to 10 to 100 h.
From the data presented in Figure 6(b)–(d), it is evident that the change in crystallographic structure of Al2Cu phase from G-P zones (155°C) to a metastable phase (190°C) to a stable phase (350°C) is the main parameter controlling the alloy performance quality. As can be seen, at each aging temperature, all points fall within a narrow circle due the progress in the formation of the precipitated phase. The broken lines in these figures show the change in the Q-level as a function of aging temperature. The width of the circle deceased from 175 MPa (155°C) to 75 MPa (190°C) to 25 MPa (350°C), representing the hardening and softening behavior of the alloy as a function of the aging temperature and time [26]. Using aging times of 2 and 100 h as reference points, the Q , UTS and %El values are presented in Table 5. As can be seen, the Q values after 2 h are more-or-less same over such a large range of aging temperatures, due to the variation in both UTS and %El. However, aging for 100 h revealed highest value at 190°C compared to 155°C (under aging) and 350°C (over aging). The Q values for test bars aged at 350°C for 100 h is the same due to the balance between UTS and %El.
\nAging temperature (°C) | \n\nQ (MPa) | \nUTS (MPa) | \n%El | \nAging time (h) | \n
---|---|---|---|---|
155 | \n432 | \n300 | \n3.5 | \n2 | \n
190 | \n432 | \n388 | \n1.5 | \n2 | \n
350 | \n406 | \n221 | \n5.3 | \n2 | \n
155 | \n340 | \n387 | \n0.9 | \n100 | \n
190 | \n470 | \n324 | \n1.5 | \n100 | \n
350 | \n406 | \n198 | \n5.3 | \n100 | \n
\nQ , UTS and %El values for alloy A after 2 and 100 h aging at different temperatures.
The heat treatment procedures followed for the alloy B are listed in Table 6. The same treatments were applied for both 25°C and 250°C tensile testing.
\nHeat treatment procedures and parameters | \n|||
---|---|---|---|
Heat treatment | \nSolution treatment | \nQuenching | \nAging | \n
SHT\n*\n\n | \n495°C for 5 h | \nWarm water (60°C) | \nNA | \n
T5 temper | \nN/A | \nN/A | \n180°C for 8 h | \n
T6 temper | \n495°C for 5 h | \nWarm water (60°C) | \n180°C for 8 h | \n
Heat treatment procedures and parameters applied to alloys investigated in stage II.
SHT, solution heat treatment.
\nFigure 7 [27] shows the DSC heating curves of the alloys in the as-cast and SHT conditions, where three explicit peaks could be detected and coded 1, 2, 3. Considering the main parameter is the precipitation of Al2Cu phase particles, thus the height of peak number 1 following SHT compared to that in the as-cast condition plays a crucial role in controlling the alloy performance after aging. In addition, it is an indication of the effectiveness of the SHT process in dissolving the initial Al2Cu phase. In Figure 7, peak # 1 after solutionizing is more or less negligible due to dissolution of most of the Al2Cu phase, as shown in Figure 4(b).
\nPortion of the DSC heating curves of as-cast and as-quenched alloy B samples [27].
The principle phases seen in alloy B are demonstrated in the optical as well as backscattered (BSE) images displayed in Figure 8(a) and (b) [27], respectively. Figure 8(a) exhibits α-Al dendrites separated by eutectic silicon colonies. The phases observed in Figure 8(b) were identified using EDS analysis and reference to the results of Hernandez-Sandoval [28] and Garza-Elizondo [29]. Selective EDS spectra identifying these phases are displayed in Figure 8(c) through Figure 8(d). The existence of Al2Cu phase in the block-like form may be attributed to the presence of Sr. in the alloy which leads to segregation of copper to localized areas [30]. The platelets of the Fe-rich β-Al5FeSi phase are easily recognized in the BSE image, surrounded by the blocky Al2Cu particles. The Mg-rich Q-phase (Al5Cu2Mg8Si6) is found growing out of the Al2Cu phase as seen in the BSE image. The absence of coarse Al3Zr precipitates [31] may be related to superheating that led to considerable dissolution of the Al3Zr phase from the master alloy during the melting process. As a result, the coarse Zr-containing phases are rarely detected since Al3Zr particles act as nucleation spots for these coarse phases. According to Garza-Elizondo [29], coarse Zr-rich particles may be nucleated on the undissolved Al3Zr particles provided by the master alloy, i.e., Al-15 wt.%Zr. In the present study, superheating the melt to 800°C would significantly reduce the numbers of Al3Zr particles in the matrix. The predicted fine zirconium trialuminide (Al3Zr) dispersoids that may be present on a nanoscale would require a high magnification BSE image to be detected.
\n(a) Optical micrograph at 200× magnification, and (b) backscattered electron image of alloy B (354 + 0.3 wt.%Zr), obtained at a low cooling rate of 0.35°C/s, showing the different phases present in the alloy; (c–g) EDS spectra corresponding to Al2Cu, (Al,Si)3(Ti,Zr), Q-Al5Mg8Cu2Si6, (Al,Si)3Zr, and β-Al5FeSi phases observed in (b) [27].
\nFigure 9(a) shows a bright-field (BF) TEM image obtained in a T6-treated sample of alloy B with the electron beam parallel to the [001] zone axis. This figure shows a high density of uniformly distributed needle-like precipitates which are oriented along <110> family of directions and aligned along the [100] planes. The length of these precipitates ranges from 50 to 150 nm, close to the reported size of θ′-Al2Cu plates (50–100 nm long) [32, 33]. Figure 9(b) displays the associated selected area electron diffraction (SAED) pattern obtained from Figure 9(a). The observable discrete diffraction maxima for the precipitates in SAED pattern indicate the presence of θ′-Al2Cu, where the streaks most probably result from the presence of fine S′-Al2CuMg particles. Computer simulation studies [34, 35, 36, 37] on the S′-phase reflections show that they are hidden within the streaks of θ′.
\n(a) Bright-field TEM image of alloy B in T6-treated condition, and (b) the selected area electron diffraction (SAED) pattern.
In the present work, Figure 10(a), the addition of ~0.3 wt.%Zr to the 354-type Al-Si-Cu-Mg cast alloy in the as-cast condition improves the ambient-temperature (25°C) strength values of the Zr-free 354 alloy (alloy A), by ~26 MPa (UTS) and 40 MPa (YS), respectively. Following SHT, the UTS and ductility values remained almost constant at ~300 MPa and ~ 6.3%, respectively, while the yield strength increased by ~33 MPa compared to alloy A. It is believed that the improved strength values of alloy B emphasizes the role of Zr addition in enhancing the ambient-temperature tensile properties through the formation of fine secondary strengthening precipitates (Al3Zr) as reported by many authors [14, 38, 39, 40]. The fact that UTS and YS in the T5 and T6 conditions are very close may be attributed to the strengthening effect of the fine dispersoids, which precipitate during the artificial aging stage of the T5, and T6 treatments as reported in Figure 9.
\n(a) Ambient, and (b, c) high temperature tensile properties of alloy B.
Tensile testing at 250°C, endured a significant softening due to the possible coarsening of the strengthening precipitates (Al2Cu) that existed during tensile testing at room temperature (Figure 11). In addition, the T5 heat treatment did not improve the elevated-temperature strength values of the as-cast alloys but reduced the alloy ductility by ~50%. However, application of the T6 heat treatment noticeably enhanced the strength values of the as-cast condition from about 175 to 225 MPa. Another parameter to consider is the effect of thermal stability. In the present work, some tensile samples were stabilized at 250°C (following the T5 and T6 aging treatment) for a lengthy period of time, i.e., 100 and 200 h. As can be seen from Figure 10(c), the stabilized T5-treated alloy B samples exhibit better strength values (UTS and YS) than those obtained in the stabilized T6-treated condition. However, the ductility values obtained after stabilization of the T5-treated samples are dramatically lower in comparison. Figure 11(a) shows Al2Cu particle size and distribution in a T6 sample stabilized at 250°C for 200 h, whereas Figure 11(b) is the corresponding EDS spectrum.
\n(a) Backscattered electron images showing the size and distribution of precipitates in the T6-treated B alloy after stabilization at 250°C for 200 h; (b) EDS spectrum corresponding to the rod-like particles in (a).
A detailed investigation of the fracture surfaces of tensile bars of alloy B were examined in the T6-treated condition, before and after stabilization for 200 h at 250°C. The T6-temper treatment was selected due to its wide use in the automotive industry. The BSE image shown in Figure 12(a) [41] shows the fracture surface of the tensile-tested alloy in the T6-treated condition. The fracture surface has a dimpled-structure throughout, which indicates the ductile nature of the fracture mode. In addition, the BSE image exhibits the precipitation of Alx(Zr,Ti)Si compound, in the form of star-like shape, as confirmed by the associated EDS spectrum in Figure 12(b). Also, cracks can be spotted in various particles of this compound, as indicated by the arrows. The higher magnification BSE image shown in Figure 11(c) reveals a cracked Alx(Zr,Ti)Si phase particle.
\nFracture surface of T6-treated alloy B: (a) BSE image showing a uniform dimple structure and cracked particles (arrowed), (b) EDS spectrum corresponding to the point of interest in (a), and (c) high magnification BSE image showing a cracked Al-Si-Ti-Zr particle (arrowed) [41].
\nFigure 13(a) [41] shows the fracture surface of the T6-treated B alloy tested at 250°C after stabilization for 200 h at the testing temperature. The dimple structure is coarser compared to that before stabilization at 250°C. This observation would explain the improved ductility of the alloy due to the softening behavior associated with the prolonged elevated-temperature exposure at 250°C. Coarsened precipitates appear in the interiors of the dimples, as indicated by the oval contours in Figure 13(a). The BSE image and the EDS spectrum shown in Figure 13(b) and (c), respectively, confirm the presence of Alx(Zr,Ti)Si phase particles.
\nFracture surface of alloy B: (a, b) BSE images of T6-treated alloy after stabilization at 250°C for 200 h showing a coarse dimpled structure, coarsened precipitates and Alx(Zr,Ti)Si particles involved in the crack initiation process, and (c) corresponding EDS spectrum of the phase of interest as shown in (b) [41].
Based on an analysis of the results presented in this article, the following conclusions may be made:
For the base 354 alloy A, solution heat treatment and artificial aging at 190°C for 2 h or at 155°C for 100 h resulted in an increase in the alloy strength by ~64% over its as-cast strength. Aging at 155 or 170°C for a long period of time offered maximum resistance to softening.
The Zr-rich intermetallic phases appear in two different forms, namely (Al,Si)2(Zr,Ti) in block-like form, and containing high level of silicon, and (Al,Si)3(Zr,Ti) in needle-like form, containing high level of aluminum.
Quality index (Q) charts constructed for alloy 354 characterize the tensile properties in terms of the heat treatment conditions applied. Minimum and maximum Q values, i.e., 259 and 459 MPa, are observed for the as-cast and solution heat-treated conditions, respectively; the yield strength shows a maximum of 345 MPa and a minimum of 80 MPa within the range of aging treatments applied.
DSC runs carried out on alloy B (354 alloy + 0.3%Zr) revealed peak patterns which included differences in peak heights—which reflected the amount of the precipitated phase, and shifts in the transformation temperature.
Melt superheating at 800°C is beneficial in terms of reducing the amount of coarse Zr-rich phases in the alloy structure, as it provides efficient dissolution of the Al3Zr phase from the master alloy during the melting process. Coarse Zr-containing phases are rarely observed due to the limited number of Al3Zr particles available to act as nucleation sites for these coarse phases.
TEM investigations confirm that the investigated alloys are strengthened primarily by θ-Al2Cu and S-Al2CuMg precipitates and their precursors, in addition to a secondary strengthening effect by precipitates in the form of Alx(Zr,Ti)Si which form following the addition of Zr.
Prolonged exposure at 250°C, resulting in coarsening of the strengthening precipitates, causes noticeable reduction in strength values, particularly the yield strength (cf. 160 and 325 MPa), and a remarkable increase in the ductility values (cf. 6.3 and 1.1%).
The strength values (UTS and YS) obtained at room temperature for the stabilized T5-treated alloy samples are comparable to those of the stabilized T6-treated condition, and higher in the case of elevated-temperature tensile testing.
The fracture surface of the T6-treated alloy B after stabilization for 1 h at 250°C reveals a dimpled-structure throughout, indicating the ductile nature of the fracture mode.
The Alx(Zr,Ti)1-xSi complex compound is observed with star-like and blocky morphologies, with cracks appearing in various particles of this compound. By increasing the stabilization time up to 200 h, coarser and deeper dimples are formed, highlighting the improved ductility of the alloy due to the softening behavior associated with the prolonged exposure at 250°C.
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