Thermal stability and Vickers microhardness of the ECAP aluminium.
\r\n\tThe most common results are dental cavities which eventually lead to pulp infection and sometimes tooth loss if left untreated. The successful endodontic treatment cane save the tooth and alveolar bone as well. Nowadays, there are many innovative methods in order to preserve the vitality of tooth without invading the pulp, and to enhance the success rate of endodontic therapy when needed, reduce the chair time and avoid post-operative symptoms. Starting from diagnosis using the most advanced X-rays, to a professional access to the pulp chamber, and root canal shaping with the very effective instruments and irrigation in order to create the ideal space for an optimal canal filling; this academic book will provide a comprehensive approach to a successful endodontic therapy of the most challenging cases, and ways of managing many severe cases in order to upgrade the life quality of the patient.
",isbn:null,printIsbn:"979-953-307-X-X",pdfIsbn:null,doi:null,price:0,priceEur:0,priceUsd:0,slug:null,numberOfPages:0,isOpenForSubmission:!1,hash:"cb2c892352c013fcad77cef2f16a1da6",bookSignature:"Ph.D. Alaa Eddin Omar Al Ostwani",publishedDate:null,coverURL:"https://cdn.intechopen.com/books/images_new/9071.jpg",keywords:"Root Canal Shaping, Ultrasonic, Reciprocating System, Rotary Files, Sodium Hypochlorite, Chlorhexidine Gel, Laser Activation, Apical Negative Pressure, Bioceramic, Gutta-percha, Vertical Compaction, MTA",numberOfDownloads:null,numberOfWosCitations:0,numberOfCrossrefCitations:0,numberOfDimensionsCitations:0,numberOfTotalCitations:0,isAvailableForWebshopOrdering:!0,dateEndFirstStepPublish:"May 7th 2019",dateEndSecondStepPublish:"September 4th 2019",dateEndThirdStepPublish:"November 3rd 2019",dateEndFourthStepPublish:"January 22nd 2020",dateEndFifthStepPublish:"March 22nd 2020",remainingDaysToSecondStep:"2 years",secondStepPassed:!0,currentStepOfPublishingProcess:5,editedByType:null,kuFlag:!1,biosketch:null,coeditorOneBiosketch:null,coeditorTwoBiosketch:null,coeditorThreeBiosketch:null,coeditorFourBiosketch:null,coeditorFiveBiosketch:null,editors:[{id:"240870",title:"Ph.D.",name:"Alaa Eddin Omar",middleName:null,surname:"Al Ostwani",slug:"alaa-eddin-omar-al-ostwani",fullName:"Alaa Eddin Omar Al Ostwani",profilePictureURL:"https://mts.intechopen.com/storage/users/240870/images/system/240870.jpeg",biography:"Dr. Al Ostwani Alaa Eddin Omar received his Master in dentistry from Damascus University in 2010, and his Ph.D. in Pediatric Dentistry from Damascus University in 2014. Dr. Al Ostwani is an assistant professor and faculty member at IUST University since 2014. \nDuring his academic experience, he has received several awards including the scientific research award from the Union of Arab Universities, the Syrian gold medal and the international gold medal for invention and creativity. Dr. Al Ostwani is a Member of the International Association of Dental Traumatology and the Syrian Society for Research and Preventive Dentistry since 2017. He is also a Member of the Reviewer Board of International Journal of Dental Medicine (IJDM), and the Indian Journal of Conservative and Endodontics since 2016.",institutionString:"International University for Science and Technology.",position:null,outsideEditionCount:0,totalCites:0,totalAuthoredChapters:"1",totalChapterViews:"0",totalEditedBooks:"1",institution:{name:"Islamic University of Science and Technology",institutionURL:null,country:{name:"India"}}}],coeditorOne:null,coeditorTwo:null,coeditorThree:null,coeditorFour:null,coeditorFive:null,topics:[{id:"16",title:"Medicine",slug:"medicine"}],chapters:null,productType:{id:"1",title:"Edited Volume",chapterContentType:"chapter",authoredCaption:"Edited by"},personalPublishingAssistant:{id:"247865",firstName:"Jasna",lastName:"Bozic",middleName:null,title:"Ms.",imageUrl:"https://mts.intechopen.com/storage/users/247865/images/7225_n.jpg",email:"jasna.b@intechopen.com",biography:"As an Author Service Manager, my responsibilities include monitoring and facilitating all publishing activities for authors and editors. From chapter submission and review to approval and revision, copyediting and design, until final publication, I work closely with authors and editors to ensure a simple and easy publishing process. I maintain constant and effective communication with authors, editors and reviewers, which allows for a level of personal support that enables contributors to fully commit and concentrate on the chapters they are writing, editing, or reviewing. I assist authors in the preparation of their full chapter submissions and track important deadlines and ensure they are met. I help to coordinate internal processes such as linguistic review, and monitor the technical aspects of the process. As an ASM I am also involved in the acquisition of editors. Whether that be identifying an exceptional author and proposing an editorship collaboration, or contacting researchers who would like the opportunity to work with IntechOpen, I establish and help manage author and editor acquisition and contact."}},relatedBooks:[{type:"book",id:"7060",title:"Gingival Disease",subtitle:"A Professional Approach for Treatment and Prevention",isOpenForSubmission:!1,hash:"b81d39988cba3a3cf746c1616912cf41",slug:"gingival-disease-a-professional-approach-for-treatment-and-prevention",bookSignature:"Alaa Eddin Omar Al Ostwani",coverURL:"https://cdn.intechopen.com/books/images_new/7060.jpg",editedByType:"Edited by",editors:[{id:"240870",title:"Ph.D.",name:"Alaa Eddin Omar",surname:"Al Ostwani",slug:"alaa-eddin-omar-al-ostwani",fullName:"Alaa Eddin Omar Al Ostwani"}],productType:{id:"1",chapterContentType:"chapter",authoredCaption:"Edited by"}},{type:"book",id:"6550",title:"Cohort Studies in Health Sciences",subtitle:null,isOpenForSubmission:!1,hash:"01df5aba4fff1a84b37a2fdafa809660",slug:"cohort-studies-in-health-sciences",bookSignature:"R. 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Venkateswarlu",coverURL:"https://cdn.intechopen.com/books/images_new/371.jpg",editedByType:"Edited by",editors:[{id:"58592",title:"Dr.",name:"Arun",surname:"Shanker",slug:"arun-shanker",fullName:"Arun Shanker"}],productType:{id:"1",chapterContentType:"chapter",authoredCaption:"Edited by"}},{type:"book",id:"72",title:"Ionic Liquids",subtitle:"Theory, Properties, New Approaches",isOpenForSubmission:!1,hash:"d94ffa3cfa10505e3b1d676d46fcd3f5",slug:"ionic-liquids-theory-properties-new-approaches",bookSignature:"Alexander Kokorin",coverURL:"https://cdn.intechopen.com/books/images_new/72.jpg",editedByType:"Edited by",editors:[{id:"19816",title:"Prof.",name:"Alexander",surname:"Kokorin",slug:"alexander-kokorin",fullName:"Alexander Kokorin"}],productType:{id:"1",chapterContentType:"chapter",authoredCaption:"Edited by"}},{type:"book",id:"314",title:"Regenerative Medicine and Tissue Engineering",subtitle:"Cells and Biomaterials",isOpenForSubmission:!1,hash:"bb67e80e480c86bb8315458012d65686",slug:"regenerative-medicine-and-tissue-engineering-cells-and-biomaterials",bookSignature:"Daniel Eberli",coverURL:"https://cdn.intechopen.com/books/images_new/314.jpg",editedByType:"Edited by",editors:[{id:"6495",title:"Dr.",name:"Daniel",surname:"Eberli",slug:"daniel-eberli",fullName:"Daniel Eberli"}],productType:{id:"1",chapterContentType:"chapter",authoredCaption:"Edited by"}},{type:"book",id:"57",title:"Physics and Applications of Graphene",subtitle:"Experiments",isOpenForSubmission:!1,hash:"0e6622a71cf4f02f45bfdd5691e1189a",slug:"physics-and-applications-of-graphene-experiments",bookSignature:"Sergey Mikhailov",coverURL:"https://cdn.intechopen.com/books/images_new/57.jpg",editedByType:"Edited by",editors:[{id:"16042",title:"Dr.",name:"Sergey",surname:"Mikhailov",slug:"sergey-mikhailov",fullName:"Sergey Mikhailov"}],productType:{id:"1",chapterContentType:"chapter",authoredCaption:"Edited by"}},{type:"book",id:"1373",title:"Ionic Liquids",subtitle:"Applications and Perspectives",isOpenForSubmission:!1,hash:"5e9ae5ae9167cde4b344e499a792c41c",slug:"ionic-liquids-applications-and-perspectives",bookSignature:"Alexander Kokorin",coverURL:"https://cdn.intechopen.com/books/images_new/1373.jpg",editedByType:"Edited by",editors:[{id:"19816",title:"Prof.",name:"Alexander",surname:"Kokorin",slug:"alexander-kokorin",fullName:"Alexander Kokorin"}],productType:{id:"1",chapterContentType:"chapter",authoredCaption:"Edited by"}}]},chapter:{item:{type:"chapter",id:"38896",title:"Equal-Channel Angular Pressing and Creep in Ultrafine-Grained Aluminium and Its Alloys",doi:"10.5772/51242",slug:"equal-channel-angular-pressing-and-creep-in-ultrafine-grained-aluminium-and-its-alloys",body:'\n\t\tCreep strength and ductility are the key creep properties of creep-resistant materials but these properties typically have opposing characteristics. Thus, materials with conventional grain sizes may be strong or ductile but there are rarely both. In this connection, recent findings of high strength and good ductility in several submicrometer metals and alloys are of special interest [1]. Reduction of the grain size of a polycrystalline material can be successfully produced through advanced synthesis processes such as the electrodeposition technique [2] and severe plastic deformation SPD [1,3-6]. Although creep is an exceptionally old area of research, above mentioned processing techniques have become available over the last two decades which provide an opportunity to expand the creep behaviour into new areas that were not feasible in earlier experiments. Creep testing of nanocrystalline (grain size d < 100 nm) and ultrafine-grained (d < 1 μm) materials is characterized by features that may be different from those documented for coarse-grained materials and thus cannot easily be compared.
\n\t\t\tProcessing through the application of severe plastic deformation (SPD) is now an accepted procedure for producing bulk ultrafine-grained materials having grain sizes in the submicrometer or nanometer range. The use of SPD enhances certain material properties through the introduction of an ultrafine-grained microstructure. The ultrafine size of the grains in the bulk materials generally leads to significantly improved properties by comparison with polycrystalline materials having conventional grain sizes of the same chemical composition. Several SPD processing techniques are currently available but the most attractive technique is equal-channel angular pressing (ECAP), where the sample is pressed through a die constrained within a channel bent through an abrupt angle [4]. There are numerous reports of the processing of various pure metals and metallic alloys by ECAP and many of these reports involve a characterization of the microstructure and an investigation of the mechanical properties at ambient temperatures. There are also several reports of the tensile properties of the as-pressed materials at elevated temperatures with a special emphasis on the potential for achieving high superplastic elongations. However, the tests at elevated temperatures are invariably conducted under conditions of constant strain rate and, by contrast, only very limited reports are available describing the creep behaviour of aluminium and some aluminium alloys. Furthermore, the results for high-purity aluminium, which are the most extensive available to date, appear anomalous because under some testing conditions of stress and temperature the measured minimum or steady-state creep rates in the pressed materials with ultrafine grain sizes where slower than in the same material in a coarse-grained unpressed condition.
\n\t\t\tThis chapter was initiated to provide basic information on the creep behaviour and microstructural characteristics of aluminium and some aluminium alloys. The chapter has arisen in connection with long-term research activity of the Advanced High Temperature Materials Group at the Institute of Physics of Materials, Academy of Sciences of the Czech Republic in Brno, Czech Republic. Thus, the objective of this chapter is to present an overview of some results of our current research in creep behaviour and a link between the microstructure and the creep properties of ultrafine-grained aluminium based alloys. Throughout the text, our results are compared with theoretical models and relevant experimental observations published in the literature.
\n\t\tProcessing by severe plastic deformation (SPD) may be defined as those metal forming procedures in which a very high strain is imposed on a bulk solid without the introduction of any significant change in the overall dimensions of the solid and leading to the production of exceptional grain refinement to that the processed bulk solids have 1000 or more grains in section [4]. Of a wide diversity of new SPD procedures, equal-channel angular pressing (ECAP) is an especially attractive processing technique. It is relatively simple procedure which can be applied to fairly large billets of many materials ranging from pure metals to precipitation-hardened alloys, intermetallics and metal-matrix composites.
\n\t\t\tThe principle of ECAP is illustrated schematically in Figure 1. For the die shown in Figure 1, the internal channel is bent through an abrupt angle, Φ, and there is an additional angle, Ψ, which represents outer arc of curvature where the two channels intersect. The sample, in the form of a rod or bar, is machined to fit within channel and the die is placed in some form of fuss so that the sample can be pressed through the die using a plunger. The nature of the imposed deformation is simple shear which occurs as the billet passes through the die. The retention of the same cross-sectional area when processing by ECAP, despite the introduction of very large strains, is the important characteristic of SPD processing and it is characteristic which distinguishes this type of processing from conventional metal-working operations such as rolling, extrusion and drawing. Since the cross-sectional area remains unchanged, the same billet may be pressed repetitively to attain exceptionally high strain.
\n\t\t\t\tPrinciple of ECAP.
Aluminium and its alloys used in this investigation were pressed using an experimental facility for ECAP installed in the Institute of Physics of Materials, Academy of Sciences of the Czech Republic (Figure 2). The die was placed on a testing machine Zwick. ECAP was conducted mostly at room temperature with a die that had internal angle 90° between two parts of the channel and an outer arc of curvature of ~ 20°, where these two parts intersect. It can be shown from first principles that these angles lead to an imposed strain of ~ 1 in each passage of the sample. The ECAP die involved the use of billets of the length of ~ 50 – 60 mm with square cross-section of 10 mm x 10 mm. The velocity of plunger was 10 mm/min.
\n\t\t\tThe use of repetitive pressing provides an opportunity to invoke different slip systems on each consecutive pass by simply rotating the samples in different ways. The four different processing routes are summarized schematically in Figure 3 [7]. In route A the sample is pressed without rotation, in route BA the sample is rotated by 90° in alternate directions between consecutive passes, in route BC the sample is rotated by 90° in the same sense (either clockwise or counter clockwise) between each pass and in route C the sample is rotated by 180° between passes. The distinction between these routes and the difference in number of ECAP passes may lead to variations both in the macroscopic distortions of the individual grains [8] and in the capability to develop a reasonably homogeneous and equiaxed ultrafine-grained microstructure.
\n\t\t\t\tAdaptation of testing ZWICK machine for ECAP pressing (a, b), and (c) sketch of ECAP die design.
Schematic of four ECAP routes for repetitive pressing.
In this work the ECAP pressing was conducted in such a way that one or repetitive pressing was conducted followed either route A, B (route BC was used only) or C. Detailed examinations of the effect of different processing routes showed that route BC leads to the most rapid evolution into an array of high-angle grain boundaries [9,10]. The result is explained by considering the shearing patterns developed in the samples during each processing route. Thus, the route BC is most probably the optimum ECAP processing route at least for the pressing of pure aluminium and its alloys [4].
\n\t\t\tDuring the last two decades it has been demonstrated that an ultrafine-grained structure of materials processed by ECAP may lead to significantly higher strength and hardness but to a reduction in the ductility [4]. In this connection after ECAP the mechanical properties were tested mostly at room temperature using a testing machine operating at a constant rate of 2.0 x 10-4s-1 of crosshead displacement.
\n\t\t\t\tTensile tests were conducted at 293 K on pure aluminium after processing by ECAP for samples after different number of ECAP passes. In limited extent mechanical tests were performed on the samples after ECAP and static annealing at 473 K [11]. In Figure 4 the tensile data are summarized as a function of the number of passes. It is apparent from these figures that a very significant increase in yield and ultimate tensile stress occurred after the first pressing. The subsequent pressing further increased yield and ultimate stress values but to a lower rate. Further, a saturation of the level of both the parameters was attained after four passes.
\n\t\t\t\t\tInfluence of different ECAP routes and different number of ECAP passes on (a) yield stress, and (b) ultimate tensile stress after static annealing.
From Figure 4 it can be also noticed that static annealing at 473 K leads to a substantial decrease in the level of yield and ultimate tensile stress values due to diffusion based recovery processes for all the ECAP processed samples. No significant differences in mechanical properties among the ECAP process routes examined were found. Further, from Figure 4 is clear that although the levels of the tensile data for ECAPed Al highly decrease with the number of ECAP passes, the stress levels after 8 passes are much higher than the stress levels in the annealing state and these differences come to more than twice. This result indicates that, when compared with the tensile behaviour of the annealed state, the flow stress is considerably improved through the application of ECAP [11,12].
\n\t\t\t\t\n\t\t\t\t\t\tFigure 5a shows Vickers microhardness plotted against the number of ECAP passes for extremely high purity aluminium (99.99%) [12]. The hardness increases up to two passes to take a maximum due to the very high dislocation density. However, subsequent passes lead to a decrease in the hardness because many of the subgrain boundaries evolve into high-angle grain boundaries. Figure 5b shows Vickers microhardness plotted against different periods of time of a static annealing at 473 K for pure (99.99%) Al processed by ECAP by two different processing routes. A pronounced decrease of microhardness with an increase of annealing time can be explained by significant grain growth and softening of pressed material during an annealing exposures [11].
\n\t\t\t\t\tHardness changes (a) with respect to number of ECAP passes, and (b) as a function of annealing time at 473 K for two different ECAP routes.
It is generally recognized that the ECAP process could produce a submicrocrystalline bulk material with a relatively uniform structure and 100% density for a wide range of materials from pure metals, solid-solution alloys, commercial alloys, to metal matrix composites [1]. However, the previously performed analysis of the data on the influence of the number of passes of equal-channel angular pressing on the elastic-plastic properties and defect structure of pure aluminium demonstrated that these characteristics of mechanical properties are substantially affected by the evolution of the nanoporosity formed during equal-channel angular pressing [13-15]. Thus, to determine the total volume of nanoporosity which could be generated by ECAP, two selected samples of pure aluminium were pressed for a total of one (specimen A1) and four (specimen A4) ECAP passes, respectively, and for comparison reasons some part of these specimens were underwent by subsequent pressurization treatment by high hydrostatic pressure [16]. The samples were investigated by small-angle X-ray scattering (SAXS) and dilatometry [13].
\n\t\t\t\t\tSome differences were found in the fractional volume of the nanopores ΔV/V when compared specimen A1 to specimen A4. The values ΔV/Vmax correspond to the as-pressed state of specimens (after ECAP only) and the values ΔV/Vmin were evaluated for the state after ECAP and subsequent pressurization which represents a rejuvenative treatment for elimination of nanopores. The evaluated values are ΔV/Vmax = 5x10-3 and ΔV/Vmin = 2.5.10-3 for specimen A1 and ΔV/Vmax = 7x10-3 and ΔV/Vmin = 3x10-3 for specimen A4, respectively. No substantial difference in the average size of the nanopores (~ 20-30nm) was found between the specimens investigated. The values ΔV/V determined by small-angle X-ray scattering and dilatometry were about the same; e.g. the fractional volume ΔV/Vmin = 3x10-3 by (SAXS) of specimen A4 agreed very well with ΔV/Vmin = 2.5x10-3 as determined by dilatometry. On the basis of the aforementioned results we can conclude that ECAP deformation achieves strongly enhanced concentration of vacancy agglomerates type defects. The effect of the spectrum of the point defects and the internal stresses on elasticity and anelasticity of ECAPed aluminium has been reported elsewhere [17].
\n\t\t\t\t\tIn recent years using a back-pressure ECAP facilities [4] has become an area of special interest. An important advantage in imposing a back-pressure may be a decrease of nanoporosity in the pressed material [18]. However, additional experiments are needed to evaluate the role of a back-pressure in elimination of nanoporosity.
\n\t\t\t\tUltrafine-grained (UFG) materials processed by ECAP differ qualitatively and quantitatively from their coarse-grained (CG) counterparts in terms of their characteristic structural parameters and thus their creep behaviour cannot be easily compared with that documented for CG materials. It is important to note in this respect that UFG materials are characterized by great extension of internal interfaces; therefore, grain boundary diffusion processes have to be involved in the formation of their structure-sensitive properties, especially at elevated temperature [19].
\n\t\t\tThe characteristics of the microstructures introduced by ECAP have been evaluated in numerous investigations [4]. However, most of these earlier investigations employed transmission electron microscopy (TEM) for determinations of the grain sizes produced by ECAP and the nature of any dislocation interactions occurring within grains. The application of modern imaging methods to the examination of microstructures in UFG materials processed by ECAP has permitted a more detailed investigation of a possible link between internal microstructures of UFG metals and alloys and their mechanical and/or creep behaviour [4]. Diffraction-based techniques for localized crystal orientation measurements, such as electron backscatter diffraction (EBSD), are of central importance today for characterizing fine-scale microstructural features [20-23].
\n\t\t\tThe new experimental technique of EBSD considerably extended the possibilities of metallography to estimate reliably the quantitative structural characteristics of materials [23]. It enables the numerical classification of boundaries separating the regions of different orientations of their lattice structure. The magnitude of the mutual misorientation can be continuously selected and thus the regions with a misorientation less than a prescribed value as well as their boundaries can be recognized. There is a vast literature devoted to the observation by EBSD and precisely defined misorientation of boundaries and the conventional grain boundary classification based on suitably polished and etched planar surfaces as observed by optical microscopy or by boundaries observed by electron microscopy and EBSD (see e.g. [24]). As can be expected, the EBSD method is more reproducible, independent of detailed etching conditions etc., and the surface area intensities are usually higher (equivalently, the mean random profile chord is smaller). In this section a division of boundaries into true subboundaries with misorientations Δ < 10°, transitional subboundaries with 10° ≤ Δ < 15°and high-angle grain boundaries with Δ ≥ 15° was made.
\n\t\t\tSuch an approach is of primary importance in the examination of materials produced by severe plastic deformation (SPD), without change of shape, producing materials with ultrafine grains (e.g. [3,5]) and considerably different properties in comparison with CG materials. The reason for this difference is to a certain degree purely geometric and consists in different grain and subgrain boundary structures, which play an important role in mechanical, thermal and other properties.
\n\t\t\tThis section describes the results of structural examinations of high purity aluminium and its selected precipitation-strengthened alloys processed by ECAP. The microstructure was revealed by TEM, SEM and EBSD and analyzed quantitatively by stereological methods. The various factors influencing the as-pressed microstructures including the total strain imposed in ECAP processing, the processing routes and the nature of materials are examined in detail.
\n\t\t\tThe aluminium used in this investigation was an extremely coarse-grained (grain size ~ 5 mm) high purity (99.99%) Al supplied in the form of rods. The rods were cut into short billets having a length of ~ 60 mm and a cross-section 10 mm x 10 mm. ECAP was conducted at room temperature using route A, BC and C. Full details on the processing have been described elsewhere [25-27].
\n\t\t\t\t\tTEM results have shown that one ECAP pass leads to a substantial reduction in the grain size (~ 1.4 μm), and the microstructure consists of parallel bands of grains oriented in the shearing direction. The microstructure is very inhomogeneous and the grain size varies from location to location. The inhomogeneous nature of the microstructure may reflect the coarse grain size (~ 5 mm) prior to ECAP. The grains subsequently evolve upon subsequent ECAP passes into a reasonably equiaxed and homogeneous microstructure with an average grain size of ~1 μm regardless of the particular ECAP routes. The microstructure is essentially homogeneous after four ECAP passes, although a tendency for grain elongation in the direction of the shear direction of the last pressing operation is retained. Figure 6 gives an example of the microstructure in the cross-section normal to the pressing direction after four subsequent ECAP passes performed in different routes. TEM micrographs in Figure 7 give an example of the microstructure in cross-section after four and eight subsequent ECAP passes by route Bc and C, respectively. The EBSD grain maps in Figure 8 indicate little dependence of the grain boundary disorientation distribution on the ECAPed Al processed by route Bc.
\n\t\t\t\t\tTEM micrographs of aluminium after four subsequent ECAP passes on route (a) A, and (b) B.
Typical microstructures and associated SAED patterns after passage through the die for (a) 4 pressings, route B and (b) 8 pressings, route C.
Grain maps for ECAPed Al after: (a) 4 passes, and (b) 8 passes by route B (EBSD).
It can be expected that the creep behaviour of the ultrafine-grained pure aluminium will critically depend on the thermal stability of the microstructure. To explore the thermal stability of ECAP processed aluminium load-less annealing was conducted at temperature of 473 K for different periods of time (i.e. at the temperature of the intended creep tests). Microscopic examination revealed that the post-ECAP annealing makes the ECAP microstructure quite unstable and a noticeable grain growth occurs at the very beginning of annealing (Table 1). Simultaneously, annealing at 473 K gives measurable change in the Vickers microhardness.
\n\t\t\t\t\t\n\t\t\t\t\t\t\t\t\t | \n\t\t\t\t\t\t\t\t\n\t\t\t\t\t\t\t\t\t | \n\t\t\t\t\t\t\t\t\n\t\t\t\t\t\t\t\t\t | \n\t\t\t\t\t\t\t||
\n\t\t\t\t\t\t\t\t\t | \n\t\t\t\t\t\t\t\t\n\t\t\t\t\t\t\t\t\t | \n\t\t\t\t\t\t\t\t\n\t\t\t\t\t\t\t\t\t | \n\t\t\t\t\t\t\t\t\n\t\t\t\t\t\t\t\t\t | \n\t\t\t\t\t\t\t|
no annealing 473 K/ 0.5 h 473 K/ 1 h 473 K/ 2 h 473 K/ 5 h 473 K/ 24 h 473 K/ 168h | \n\t\t\t\t\t\t\t\t0.9 6.6 7.9 7.3 7.3 12.2 13.4 | \n\t\t\t\t\t\t\t\t37 27 23 23 21 19 18 | \n\t\t\t\t\t\t\t\t0.9 4.5 4.8 4.8 5.3 5.0 10.4 | \n\t\t\t\t\t\t\t\t38 32 32 27 27 23 21 | \n\t\t\t\t\t\t\t
Thermal stability and Vickers microhardness of the ECAP aluminium.
In evaluating the microstructure characteristics of ultrafine-grained materials processed by ECAP at elevated and high temperatures, it is very important to recognize that these ultrafine-grained microstructures are frequently unstable at these temperatures as it was just demonstrated by the above mentioned results of thermal instability of pressed pure aluminium. However, it is often feasible to retain an array of ultrafine grains even at very high temperatures by using materials containing second phases or arrays of precipitates. This was a reason why two precipitation-strengthened aluminium alloys were used in this investigation.
\n\t\t\t\t\tIt has been shown that addition to aluminium alloys of even very small amounts of Sc (typically, ~ 0.2wt.%) strongly improves the microstructures of the alloys and their mechanical properties so that these alloys are suitable for use in engineering applications [28]. Scandium additions of ~ 0.2wt.%Sc to pure aluminium are sufficient to more or less retain a small grain size at elevated temperatures [29]. Further, some reports have demonstrated that it is possible to achieve high ductilities in Al-Mg-Sc alloys by using ECAP to introduce an exceptionally small grain size [30]. The creep behaviour of conventional Al-Mg alloys is extensively described in the literature. The synergy of solid-solution strengthening and precipitate strengthening has, however, not been extensively studied at elevated and high temperatures [31]. Very little information is available at present on the creep properties of ultrafine-grained Al-Sc and Al-Mg-Sc alloys [32-38]. Accordingly, the present investigation was initiated to provide a more complex information on the creep behaviour of these aluminium alloys in their ultrafine-grained states.
\n\t\t\t\t\tAn Al-0.2wt.%Sc alloy was produced by diluting an Al-2.0wt.%Sc master alloy with 99.99wt.% pure aluminium. The resulting ingots were subjected to a homogenization and grain-coarsening treatment at 893 K for 12 hours and then aged in air at 623 K for 1 hour. In the as-fabricated condition, the extremely coarse grain size was measured as ~ 8 mm. The ECAP was conducted at the Institute of Physics of Materials AS CR Brno, Czech Republic, using the same die and procedure as it was reported earlier for pure aluminium (i.e. up to a total 8 ECAP passes at room temperature). The details concerning an Al-0.2wt.%Sc alloy have been reported elsewhere [33-35]. The ternary Al-Mg-Sc alloy was fabricated at the Department of Materials Science and Engineering, Faculty of Engineering, Kyushu University, Fukuoka, Japan. The alloy contained 3wt.%Mg and 0.2wt.%Sc and it was prepared from 99.99% purity Al, 99.999% purity Sc and 99.9% purity Mg. Full details on the fabrication procedure are given elsewhere [32] but, briefly, the alloy was cast, homogenized in air for 24 h at 753 K and solution treated for 1 h at 883 K. In the as-fabricated condition, the grain size was about 200 μm. Again, the ECAP was conducted using a solid die that had 90° angle between the die channels and each sample was pressed at room temperature repetitively for a total of eight passes by route BC.
\n\t\t\t\t\tMicrostructure in the Al-0.2wt.%Sc alloy: (a) and (b) after ECAP (BC, 8 passes) and annealing for 1 h at 623 K, (c) and (d) after creep at 473 K.
\n\t\t\t\t\t\tFigures 9a,b and 10a,b show the microstructure of Al-0.2wt.%Sc and Al-3wt.%Mg-0.2wt.%Sc alloys in their as-pressed states. Experiments on Al-0.2wt.%Sc and Al-3wt.%Mg-0.2wt.%Sc alloys revealed that processing by ECAP reduced the grain size to ~ 0.4 μm and subsequent annealing at 623 K and 1 h and creep testing gave the grain sizes ~ 0.9 μm for an Al-0.2wt.%Sc alloy and ~ 1.5 μm for an Al-3wt.%Mg-0.2wt.%Sc alloy, respectively. Figures 9c,d and 10c,d give examples of the microstructure of the alloys in the longitudinal sections parallel to the pressing direction after creep exposures at 473 K. As will be shown later on no substantial difference in the relative fractions of high-angle (θ > 15°) grain boundary population after ECAP was found between the alloys investigated. These fractions were slightly increased during creep exposure up to an average value ~ 70%.
\n\t\t\t\t\tMicrostructure in the Al-3wt.%Mg-0.2wt.%Sc alloy: (a) and (b) after ECAP (BC, 8 passes) and annealing for 1 h at 623 K, (c) and (d) after creep at 473 K.
TEM micrograph showing the presence of coherent Al3Sc precipitates in unpressed sample, and (b) precipitate size distribution in unpressed sample.
\n\t\t\t\t\t\tFigures 9b,d and 10b,d exhibit TEM micrographs, which demonstrate the presence of coherent Al3Sc precipitates within the matrices of both alloys. Al3Sc precipitates are indicated by a coherency strain contrast [39]. A mean size of precipitates was ~ 5 nm after creep testing of the Al-3wt.%Mg-0.2wt.%Sc alloy and a mean size of precipitates slightly large was recorded for the Al-0.2wt.%Sc alloy (~ 6 nm [35]). Figure 10d shows dislocation microstructure observed after creep exposure of ECAPed Al-3wt.%Mg-0.2wt.%Sc. The dislocation pairs present in both alloys containing the smallest precipitate radii are very frequent. For larger precipitates the dislocations are pinned efficiently by Al3Sc precipitates as climbing becomes slower. Figure 11a exhibits TEM micrograph of an Al-3wt.%Mg-0.2 wt.%Sc alloy showing the presence of coherent Al3Sc precipitates in an unpressed sample, and Figure 11b presents precipitate size distribution in an unpressed alloy.
\n\t\t\t\tIt can be expected that the creep behaviour of the UFG material will be influenced critically upon the subsequent thermal stability of its microstructure. To explore this effect microscopic examination of grain size change in pure aluminium during creep exposure at 473 K and 15 MPa were performed. It is important to note that each creep specimen was heated to the testing temperature in the furnace of the creep testing machine over a period of ~ 2h and then held at the testing temperature for further ~ 2h in order to reach thermal equilibrium. Consequently, the microstructure characteristics of the ECAP material at the onset of the creep testing were similar to that shown in Table 1. No substantial coarsening of grains has been observed during creep exposure at 473 K (see Table 2).
\n\t\t\t\t\t\n\t\t\t\t\t\t\t\t\t | \n\t\t\t\t\t\t\t\t\n\t\t\t\t\t\t\t\t\t | \n\t\t\t\t\t\t\t\t\n\t\t\t\t\t\t\t\t\t | \n\t\t\t\t\t\t\t\t\n\t\t\t\t\t\t\t\t\t | \n\t\t\t\t\t\t\t
A4 | \n\t\t\t\t\t\t\t\troute A, 4 passes | \n\t\t\t\t\t\t\t\t6.4 | \n\t\t\t\t\t\t\t\t79 | \n\t\t\t\t\t\t\t
A8 | \n\t\t\t\t\t\t\t\troute A, 8 passes | \n\t\t\t\t\t\t\t\t7.0 | \n\t\t\t\t\t\t\t\t26 | \n\t\t\t\t\t\t\t
A12 | \n\t\t\t\t\t\t\t\troute A, 12 passes | \n\t\t\t\t\t\t\t\t6.7 | \n\t\t\t\t\t\t\t\t17 | \n\t\t\t\t\t\t\t
B4 | \n\t\t\t\t\t\t\t\troute B, 4 passes | \n\t\t\t\t\t\t\t\t8.7 | \n\t\t\t\t\t\t\t\t62 | \n\t\t\t\t\t\t\t
B8 | \n\t\t\t\t\t\t\t\troute B, 8 passes | \n\t\t\t\t\t\t\t\t7.2 | \n\t\t\t\t\t\t\t\t60 | \n\t\t\t\t\t\t\t
B12 | \n\t\t\t\t\t\t\t\troute B, 12 passes | \n\t\t\t\t\t\t\t\t8.8 | \n\t\t\t\t\t\t\t\t39 | \n\t\t\t\t\t\t\t
Grain size of the ECAP material after creep at 473 K and 15 MPa.
TEM observations were used also to established details of microstructure evolution during creep. The micrographs in Figure 12a,b illustrate a dislocation substructure inside the grains. The dislocation lines were wavy and occasionally tangled with each other. It is know that large grains in UFG materials contain dislocations while grains smaller than a certain size are dislocation free [3,6]. EBSD measurements were taken to determine the grain boundary misorientation and the value of relative fraction of a high-angle grain boundary (θ > 15°) population (for details see 3.3.2.).
\n\t\t\t\t\tTEM micrographs from the longitudinal section of an aluminium processed by ECAP route Bc (a) after 1 ECAP pass and creep, (b) after 8 ECAP passes and creep. Creep at 473 K and 15 MPa.
For comparison reasons, some results of microstructural changes in Al-0.2wt.%Sc and Al-3wt.%Mg-0.2wt.%Sc alloys during creep were presented earlier in 3.1.2. It was found that creep exposures of an Al-0.2wt.%Sc alloy at 473 K and 20 MPa caused the changes in (sub)grain sizes initially resulting from ECAP pressing. Figure 13a shows the microstructure after 8 ECAP passes and subsequent creep exposure. TEM analysis revealed that the average (sub)grain size increases from ~ 0.4 um to ~ 1.3 μm after creep exposure. The (sub)grain growth was effected by presence of coherent Al3Sc precipitates (Figure 13b) which to some extent pinned the boundaries against their migration and restricted the movement of dislocation.
\n\t\t\t\t\tMicrostructure of Al-0.2wt.%Sc alloy after 8 ECAP passes and subsequent creep exposure at 473 K and 20 MPa. (a) microstructure, and (b) precipitates Al3Sc.
The EBSD data indicate that the number of high angle-boundaries (θ>15°) measured in the specimens after ECAP and subsequent creep exposure is strongly dependent on the number of ECAP passes. The number of high-angle grain boundaries is increasing with increasing number of ECAP passes from approximately 2% in the specimen after 1 ECAP pass and subsequent creep to ~ 70% in the specimen after 8 ECAP passes and subsequent creep. It was reported [4] that the grain boundary sliding can occur in UFG materials at elevated temperatures. Thus we can suppose that changes in the number of high-angle grain boundaries in the microstructure of ECAPed materials during creep tests can affect their creep behaviour by increasing the contribution of grain boundary sliding to the total creep strain [27].
\n\t\t\t\t\tThe EBSD analyses were performed on the several places of the gauge length of creep specimen after ECAP and subsequent creep revealed scatter in the number of high angle grain boundaries (HAGBs). In the Figure 14 the minimal and maximal measured values of the number of HAGBs are plotted. The inspection of Figure 14 shows that the scatter in HAGBs can be particularly expected after creep tests in the specimens with lower number of ECAP passes. The heterogeneous distribution of HAGB can probably influence the homogeneity of grain boundary sliding. In the areas with the higher number of HAGBs the grain boundary sliding will be more intensive than in the surrounding areas [8].
\n\t\t\t\t\tFraction of high angle grain boundaries as a function of the number of ECAP passes in the Al-0.2wt.%Sc alloy.
The investigation of the unetched surfaces of the specimens after 2-8 ECAP passes and after creep exposure revealed the appearance of mesoscopic shear bands [14,15,35, 40-42] lying near to the shear plane of the last ECAP pass (Figure 15). On the surface of specimens the mesoscopic shear bands were particularly observed near the fracture region and their frequency decreased rapidly with increasing distance from the fracture. On the specimen surface after 8 ECAP passes the mesoscopic shear bands already covered almost the whole gauge length. It was found that the width of the bands decreases with increasing number of ECAP passing and after 8 ECAP passes the average width of the bands was ~ 35 µm as it is shown in Figure 15. The analyses of microstructure on the interfaces of the bands found that in the vicinity of these interfaces high heterogeneity in the distribution of HAGBs can be observed (Figure 15). The formation of the mesoscopic shear band can be related to inhomogeneity of microstructure of ECAPed alloy after creep exposure. Examination by EBSD revealed that the microstructure of mesoscopic shear bands is created by high-angle grain boundaries (Figure 15 and 16).
\n\t\t\t\t\tThe heterogeneous distribution of HAGBs (red coloured) in the sample of an Al-0.2wt.%Sc alloy after 4 ECAP passes and creep at 473 K and 20 MPa.
Appearance of the microstructure in the Al-0.2wt%Sc alloy after 2 ECAP passes and subsequent creep at 473 K. Tensile axis is horizontal, SEM.
The processing technique used to obtain the UFG microstructure should strongly influence the creep properties of the material. This is primarily due to difference in microstructure as distribution of grains, subgrains, dislocation density and boundary character. The grain boundary character is usually quantified using the misorientation angle θ across grain boundaries, with high and low angle grain boundaries defined as θ ≥ 15° an 2° < θ < 15°, respectively. Electron back scatter diffractions (EBSD) mapping has been used to quantitatively characterize boundaries in UFG materials [10,20,21,23,24]. Although the boundary spacing saturates after the first few ECAP passes, the fraction of high angle boundaries continues to increase with increasing ECAP passes [27,34,43]. In addition to grain size determination, there are a number of important microstructural parameters evaluated from EBSD but not available from conventional methods of grain characterization in particulars parameters relating to the grain orientations and boundary characters [44 - 46]. The following text describes representative results of quantitative characterization of UFG microstructure.
\n\t\t\t\tAt each examined specimen there were made three mutually perpendicular planar metallographic sections denoted as XY, XZ (longitudinal sections) and YZ (transverse section), where X, Y and Z are the axes of the Cartesian coordinate system with X along the last pressing direction and Z perpendicular to the bottom of the channel. The technique of automated EBSD in the scanning electron microscope was used for quantitative metallography. Four ranges of the boundary misorientation ∆ were selected; 2° ≤ ∆, 5° ≤ ∆, 10° ≤ ∆ and 15° ≤ ∆. Then standard intercept counting [45] resulting in the mean number NL of profile chords per unit length of the examined test lines was carried out. In each specimen, six systematically selected directions of the test lines in each section were examined. The mean boundary areas unit volume were then estimated by the stereological relation SV = [2NL]. Another important feature of the grain boundary structure is its inhomogeneity. The dispersion of grain profile areas can be qualified by the coefficient of variation CVa of the grain profile areas in a plane CVa = \n\t\t\t\t\t\t\t
There are numerous reports of the processing of various pure metals and metallic alloys by ECAP and many of these reports involve a detailed characterization of the microstructure. These results are summarized in recent reviews [3,4]. However, information seldom is reported on the percentage of high angle grain boundaries (HAGB´s), an important parameter in the comparison of plasticity of different processing routes and materials [50]. It can be expected that samples with different distributions of misorientation across the grain boundaries will deform differently. Further, to provide information on the optimum microstructure of UFG materials we need to use an additional quantitative microstructural parameter other than just the average grain size critical for the creep behaviour and properties [51]. Such parameter could be a coefficient of profile CVa as a measure of homogeneity of materials microstructure [48].
\n\t\t\t\t\tHence, the grain and subgrain structure of the creep specimens was revealed by means of EBSD and characterized by the coefficient of variation CVa of the profile areas. Four ranges of the boundary misorientation ∆ between adjacent pixels were selected for examination using EBSD, which correspond namely to subboundaries, transitive and high angle grain boundaries within 2° ≤ ∆ and 5° ≤ ∆, transitive and high angle grain boundaries for 10° ≤ ∆, and mostly grains with HAGB´s for ∆ ≥ 15°. Selected examples of images of XZ sections produced by EBSD of an Al-0.2wt.%Sc are shown in Figure 17.
\n\t\t\t\t\tSelected examples of EBSD grain maps of an Al-0.2wt.%Sc alloys after ECAP constructed for different grain misorientations: (a) subboundaries ∆ ≥ 2°, (b) transitive subboundaries and high angle boundaries ∆ ≥ 10°, and, (c) high-angle grain boundaries ∆ ≥ 15°.
The fraction of high-angle grain boundaries in the crept samples as a function of the number of ECAP passes.
It was generally observed that with the increasing number of ECAP passes N, a considerable amount of subgrain boundaries was gradually transformed to HAGB´s as shown in Figure 18. At the same time, the local homogeneity of structure as characterized by the values of CVa with the increasing number N is gradually improved as demonstrates by Figure 19. Figure 19 also shows the microscopic appearance of the specimens of an Al-0.2wt.%Sc alloy crept under the same loading conditions (473 K, 20 MPa) but processed by different numbers of ECAP passes N. The values of CVa as high as 10 at N = 2, 1 ≤ CVa< 2 at N = 4, and 0.55 ≤ CVa ≤ 1 at N = 8 were found. Extremely high value CVa at N = 2 demonstrates very high inhomogeneity of a mixture of subgrain and grain structures. The value of CVa in very homogeneous grain systems should not exceed the value of 1 [23]. It should be noted that extremely high values of the coefficient of variation CVa is a natural consequence of the short as well as long-range inhomogeneity of microstructure.
\n\t\t\t\t\tGrain maps of an Al-0.2wt.%Sc alloy processed by different number N of ECAP passes and crept at 473 K and 20 MPa, and corresponding parameter CVa: (a) 2 passes, CVa>> 2, (b) 4 passes, CVa< 2 and, (c) 8 passes, 0.55 ≤ CVa< 1.
The substantial grain coarsening especially in the case of pure metals came up during the creep exposures depending on stress and temperature thus manifesting the thermal instability of ultrafine-grained microstructure [4,52,53]. It is clear that in pure aluminium the grains grow rapidly at elevated temperatures because there are no precipitates within the crystalline lattice to restrict the movement of the grain boundaries by a “pinning effect”. By contrast, submicrometer grains may be retained to relatively high temperatures in materials containing a distribution of fine precipitates as in the case of an Al-0.2%Sc alloys containing Al3Sc precipitates [29]. The quantitative characterization of the inhomogeneity of the boundary structure is shown in Figure 20. It should be stressed that the values of the coefficient of profile area CVa as a measure of structure homogeneity strongly depend on the chosen ranges of misorientation ∆ in EBSD analysis. Whereas the fraction of the subboundaries (low-angle grain boundaries) are dominating for ∆ ≥ 2° (Figure 21a), the fractions of high-angle grain boundaries (θ ≥ 15°) confirm their high share for the range of ∆ ≥ 15° (Figure 21b). Detailed inspection of Figure 20b shows a strong dependence of CVa after creep on the number of ECAP passes. Substantial decrease of the values CVa 2° for the precipitation-strengthened Al-0.2Sc alloy with increasing number of passes N for ∆ ≥ 2° may be connected with the more rapid evolution boundaries having misorientation angles θ > 15° (Figure 20a).
\n\t\t\t\t\tCoefficient of profile area CVa as a measure of homogeneity: 0.55 ≤ CVa < 1 (homogeneous system), and CVa >> 2 (multimodal grain size distribution). The chosen ranges of misorientation ∆ in EBSD analysis: (a) ∆ ≥ 2°, (b) ∆ ≥ 15°.
Distribution of boundaries with different misorientation θ for an Al-0.2wt.%Sc alloy analysed in
The mechanical properties of bulk ultrafine-grained (UFG) materials at elevated and/or high temperatures are a new and important area of research [4]. However, there have been only a few investigations on the creep behaviour of bulk UFG materials processed by equal-channel angular pressing (ECAP) [43,54]. By comparison with the unpressed (coarse-grained) state, processing by ECAP may lead to considerable changes in the creep properties in bulk UFG materials including a decrease and/or an increase [8] in the minimum creep rate and the creep life. It is important to note that these trends may be noticeably dependent on the number of ECAP passes. Many investigations concerned with the identification of creep mechanisms have been undertaking using coarse-grained pure aluminium [55] and commercial aluminium alloys [56]. However, it is logical to expect that the mechanism of hardening/softening observed in the aluminium processed by ECAP may be different from that observed in the coarse-grained material. Consequently, it cannot be excluded that creep in the ECAP aluminium and its alloys is controlled by different creep mechanism(s) than that in the coarse material. Thus, neither phenomenological nor microscopic aspects of the creep behaviour of materials processed by ECAP have been understand sufficiently as yet.
\n\t\t\tThis section reports a series of creep experiments that were conducted on specimens of pure aluminium and its Al-0.2wt.%Sc and Al-2wt.%Mg-0.2wt.%Sc alloys processed by equal-channel angular pressing. For comparison purposes, some creep tests were performed also on the unpressed materials. Creep tests were performed both in tension and compression.
\n\t\t\tAs discussed in more detail in Section 2.2. four distinct pressing routes have been identified (Figure 3). The ECAP processing was conducted by one or repetitive passes following either route A, B (route BC was used only) or C. Figure 22 contains a complete record of the main creep parameters for the ECAP specimens of Al after creep testing in tension at 473 K and 15 MPa (each point represents the average results of two to three individual creep tests at the same loading conditions). Inspection of Figure 22 shows there are not very significant differences in creep properties of specimens prepared by the various ECAP processing routes. All three processing routes produce a significant increase in the minimum creep rate through the first four passes and a slight increase during subsequent pressing (Figure 22a). By contrast, the time to fracture (creep life) dramatically drops through four passes and then there is no significant differences among the number of following passes – Figure 22b.
\n\t\t\t\tRecently, attention has been given to effectiveness of the various ECAP routes in producing grain refinement in aluminium [10]. It has been demonstrated that ECAP is capable of producing refined structures with large fractions of high-angle boundaries [10] although the mechanisms involved in the formation of fine grains and high-angle boundaries in the deformation microstructure remain to be clarified. In this work microstructural investigation all routes examined indicated little differences in the grain size produced via the various ECAP routes. With increasing number of ECAP passes this difference decreases. Further, there was little apparent dependence of the misorientation on the various process route for an ECAP die having an internal angle equals to 90°. The misorientation data confirmed that repetitive pressing results in a progressive increase in the fraction of high-angle grain boundaries (Figure 18).
\n\t\t\t\tIn related work, Sklenicka et al. [11,57] carried out an extensive creep testing on pure aluminium processed by various ECAP routes. It was found that processing route had a little apparent effect on the creep behaviour of a pressed aluminium. However, the effect of grain growth during creep may tend to obscure the effect of different processing routes and the creep experiments are probably not a sufficiently refined procedure for picking up these rather small differences in the creep behaviour.
\n\t\t\t\tInfluence of different ECAP routes and different number of ECAP passes on (a) creep rate, and (b) time to fracture.
Creep tests were performed on the as-pressed specimens both in tension and compression in the temperature interval from 423 to 523 K under an applied stress range between 10 and 25 MPa. The subsequent ECAP passes were performed by route BC (see part 2.2.) up to 12 passes.
\n\t\t\t\tRepresentative creep curves are shown in Figures 23 and 24. All of these plots were obtained at temperature of 473 K (~ 0.5 Tm) under an initial applied uniaxial tensile or compression stress of 15 MPa. The creep tests in tension were run up to the final fracture of the creep specimens, whereas the creep tests in compression were interrupted at a true strain of about ~ 0.35.
\n\t\t\t\t\tStandard ε vs t creep curves in Figures 23a and 24a can be easily replotted in the form of the instantaneous strain rate dε/dt versus strain as shown in Figures 23b and 24b. As demonstrated by figures, significant differences were found in the creep behaviour of the ECAP material when compared to its coarse-grained counterpart. First, the ECAP materials exhibits markedly longer creep life (Figure 23a) or markedly longer duration of creep exposure to obtain a strain of ~ 0.35 (Figure 24b) than coarse grained aluminium. Second, the minimum creep rate for the ECAP material is about one to two orders of magnitude less than that of coarse-grained material. Third, the shapes of tensile creep curves for the ECAP material after high number of pressing differ considerably from the tests conducted at small number of the ECAP passes by the extent of individual stages of creep.
\n\t\t\t\t\tStandard creep and creep rate versus strain curves for unpressed state and various number of ECAP passes via route Bc (creep in tension up to fracture).
Standard creep and creep rate versus strain curves for unpressed state and various number ECAP passes via route Bc (creep in compression up to strain ~ 0.35).
The difference in the minimum creep rate for the ECAP material and unpressed state consistently decreases with increasing number of ECAP passes (Figures 23b and 24b). An additional difference is illustrated by Figure 25a, which shows the variation of the minimum creep rate with the applied stress for the ECAP specimens after 8 passes. The results demonstrate that at high stresses the minimum compressive creep rate of the ECAP material may be up to one order of magnitude lower than that of the unpressed material, although this difference decreases with decreasing applied stress and becomes negligible at 10 MPa. The observed values of the stress exponent n = (\n\t\t\t\t\t\t\t
To determine the apparent activation energy for creep Qc, the minimum creep rate was measured in the temperature interval from 423 to 523 K and at two tensile applied stresses 15 and 20 MPa, respectively. The activation energy for creep Qc is defined as
\n\t\t\t\t\tThus, the activation energy QC can be derived from the slope of log dε/dt versus 1/T plots shown in Figure 25b. A value of the apparent activation energy, QC, was determined by the least square methods. The QC is stress dependent and equals to 129.7 ± 16 and 110.9 ± 9 kJ/mol for stresses 20 and 15 MPa, respectively.
\n\t\t\t\t\tDependence of minimum creep rate for unpressed state and 8 ECAP passes on: (a) applied stress, (b) testing temperature at two levels of stress.
The total creep strain ε generally consists of following contributions: the strains caused by dislocation glide and nonconservative motion of dislocations, grain boundary sliding (GBS) [58-60], stress directed diffusion of vacancies and by intergranular void nucleation and growth, respectively. However, it should be noted that not all of the above processes operating are independent of each other, as frequently assumed. A possible explanation for the occurrence of intensive GBS in UFG materials is that diffusion is more rapid in ECAP processed materials with highly non-equilibrium grain boundaries [4, 19, 61,62]. Accordingly, it appears that GBS is easier in these UFG materials.
\n\t\t\t\t\tThe amount of grain boundary sliding (GBS) was determined by measuring the surface offsets produced at the intersections of grain boundaries with marker lines transverse to the stress axis [27,58]. Figure 26 shows one clear example of the occurrence of grain boundary sliding in creep of the ECAP aluminium. Longitudinal displacements of the marker lines, u, due to GBS, together with the fraction of boundaries, κS, with observable GBS, were measured using SEM.
\n\t\t\t\t\tGrain boundary sliding was measured on the surfaces of the tensile specimens crept up to a predetermined strain ε ≈ 0.15. Scanning electron microscopy made it possible to detect GBS characterized by u ≥ 0.1μm. However, GBS was not observed at all grain boundaries; that is why the relative frequency of sliding boundaries κs was determined. Then the strain component εgb due to GBS is expressed as [27,58]:
\n\t\t\t\t\twhere the mean grain size \n\t\t\t\t\t\t\t
\n\t\t\t\t\t\t\t\t\t | \n\t\t\t\t\t\t\t\t\n\t\t\t\t\t\t\t\t\t | \n\t\t\t\t\t\t\t\t\n\t\t\t\t\t\t\t\t\t | \n\t\t\t\t\t\t\t\t\n\t\t\t\t\t\t\t\t\t | \n\t\t\t\t\t\t\t\t\n\t\t\t\t\t\t\t\t\t | \n\t\t\t\t\t\t\t\t\n\t\t\t\t\t\t\t\t\t | \n\t\t\t\t\t\t\t
1 | \n\t\t\t\t\t\t\t\t0.51 | \n\t\t\t\t\t\t\t\t0.80 | \n\t\t\t\t\t\t\t\t14.9 | \n\t\t\t\t\t\t\t\t3.15 | \n\t\t\t\t\t\t\t\t21.0 | \n\t\t\t\t\t\t\t
2 | \n\t\t\t\t\t\t\t\t0.48 | \n\t\t\t\t\t\t\t\t0.83 | \n\t\t\t\t\t\t\t\t12.7 | \n\t\t\t\t\t\t\t\t3.60 | \n\t\t\t\t\t\t\t\t24.0 | \n\t\t\t\t\t\t\t
4 | \n\t\t\t\t\t\t\t\t0.55 | \n\t\t\t\t\t\t\t\t0.93 | \n\t\t\t\t\t\t\t\t12.2 | \n\t\t\t\t\t\t\t\t4.80 | \n\t\t\t\t\t\t\t\t32.0 | \n\t\t\t\t\t\t\t
8 | \n\t\t\t\t\t\t\t\t0.49 | \n\t\t\t\t\t\t\t\t0.91 | \n\t\t\t\t\t\t\t\t10.8 | \n\t\t\t\t\t\t\t\t4.70 | \n\t\t\t\t\t\t\t\t31.0 | \n\t\t\t\t\t\t\t
12 | \n\t\t\t\t\t\t\t\t0.52 | \n\t\t\t\t\t\t\t\t0.92 | \n\t\t\t\t\t\t\t\t11.0 | \n\t\t\t\t\t\t\t\t5.00 | \n\t\t\t\t\t\t\t\t33.0 | \n\t\t\t\t\t\t\t
Summary of GBS measurements (ε \n\t\t\t\t\t\t\t\t\t
Example of grain boundary sliding in the ECAPed aluminium (route Bc, 8 passes) after creep testing at 473 K and 15 MPa. Tensile stress axis is horizontal.
The mechanisms controlling the creep properties of pure metals have been usually identified from the dependence of the minimum and/or steady-state creep rate \n\t\t\t\t\t\t\t
where Qc is the activation energy for creep. With this approach, the fact that n, p and Qc are themselves functions of stress, temperature and grain size is conventionally explained by assuming that different mechanisms, each associated with different values of n, p and Qc, control the creep characteristics in different stress/temperature regimes. In turn, the dominant mechanisms under specific test conditions are then generally determined by comparing experimentally determined values of n, p and Qc with the values predicted theoretically for different creep mechanisms.
\n\t\t\t\t\tIn Figure 27, the minimum creep rates \n\t\t\t\t\t\t\t
The broken lines in Figure 27 denote the model predictions of the theoretical creep rates according models for various creep deformation mechanisms, namely for superplastic flow, Nabarro-Herring [59, 63,64] and Coble [59,65] diffusion creep and power-law creep by dislocation climb and glide processes. It should be stressed that the phenomenon of dislocation diffusion is not well understood on the fundamental level at present. The theoretical creep rates were calculated from equation (3) using the material data presented in Table 4.
\n\t\t\t\t\t\n\t\t\t\t\t\t\t\t\t | \n\t\t\t\t\t\t\t||||||||
(1), (1*) | \n\t\t\t\t\t\t\t\tsuperplastic flow | \n\t\t\t\t\t\t\t\t2 | \n\t\t\t\t\t\t\t\t2 | \n\t\t\t\t\t\t\t\t10 | \n\t\t\t\t\t\t\t\tDGB\n\t\t\t\t\t\t\t\t | \n\t\t\t\t\t\t\t\t5.9x10-14\n\t\t\t\t\t\t\t\t | \n\t\t\t\t\t\t\t\t86 | \n\t\t\t\t\t\t\t\t[66] | \n\t\t\t\t\t\t\t
(2), (2*) | \n\t\t\t\t\t\t\t\tNabarro – Herring creep | \n\t\t\t\t\t\t\t\t1 | \n\t\t\t\t\t\t\t\t2 | \n\t\t\t\t\t\t\t\t28 | \n\t\t\t\t\t\t\t\tDL\n\t\t\t\t\t\t\t\t | \n\t\t\t\t\t\t\t\t2.72x10-20\n\t\t\t\t\t\t\t\t | \n\t\t\t\t\t\t\t\t143.4 | \n\t\t\t\t\t\t\t\t[66] | \n\t\t\t\t\t\t\t
(3), (3*) | \n\t\t\t\t\t\t\t\tCoble creep | \n\t\t\t\t\t\t\t\t1 | \n\t\t\t\t\t\t\t\t3 | \n\t\t\t\t\t\t\t\t62 | \n\t\t\t\t\t\t\t\tDGB\n\t\t\t\t\t\t\t\t | \n\t\t\t\t\t\t\t\t5.9x10-14\n\t\t\t\t\t\t\t\t | \n\t\t\t\t\t\t\t\t86 | \n\t\t\t\t\t\t\t\t[66] [67] | \n\t\t\t\t\t\t\t
(4) | \n\t\t\t\t\t\t\t\tdislocation climb and glide | \n\t\t\t\t\t\t\t\t5 | \n\t\t\t\t\t\t\t\t0 | \n\t\t\t\t\t\t\t\t103\n\t\t\t\t\t\t\t\t | \n\t\t\t\t\t\t\t\tD*\n\t\t\t\t\t\t\t\t | \n\t\t\t\t\t\t\t\t1.9x10-14\n\t\t\t\t\t\t\t\t | \n\t\t\t\t\t\t\t\t124 | \n\t\t\t\t\t\t\t\t[66] | \n\t\t\t\t\t\t\t
Creep mechanisms and material data.
D* is the effective diffusion coefficient which incorporates contribution from both lattice and grain boundary diffusion, Dgb and DL are the grain boundary and lattice diffusion coefficients, respectively.
It should be noted that grain growth occurs easily at the elevated temperatures used in creep experiments of pure metals. Indeed, the occurrence of significant grain growth in creep tests conducted on high-purity aluminium processed by ECAP at room temperature was observed [27,36]. Accordingly, two sets of the predicted theoretical rates were calculated in this analysis for both states of materials using the measured grain size after ECAP processing (dECAP) and after subsequent creep exposure (dCREEP) [27]. The theoretical rates using dECAP will be marked by simple number (e.g. 1) while a numbering with asterisk (e.g. 1*) will be used for the rates corresponding to dCREEP in Figure 27.
\n\t\t\t\t\tExperimentally determined and theoretically predicted the stress dependences of the minimum creep rates for various creep mechanisms in aluminium.
\n\t\t\t\t\t\tFigure 27 demonstrates that at high applied stresses the experimentally determined minimum creep rate of the ECAP aluminium may be up to two orders of magnitude lower than that of the unpressed material, although this difference decreases with decreasing applied stress and becomes nearly negligible at 10 MPa. The predictions show that under the creep loading conditions investigated Nabarro-Herring and Coble diffusion creep and superplastic flow are too slow to account for the creep deformation considering a significant grain growth in the pressed materials. Also shown in Figure 27 are the predicted theoretical creep rates for ultrafine-grained states (dECAP) after ECAP which are within two to five orders of magnitude faster than that for the creep of coarsened materials. However, such predictions are not correct a priori due to thermal instability of microstructure of the pressed materials. Inspection of Figure 27 shows that for the pressed material there is an excellent agreement between the experimental datum points and the predicted creep behaviour based on dislocation climb and glide. Further, for n ≥ 4 creep is known to occur by diffusion-controlled movement of dislocations within grains and/or along grain boundaries (grain boundary sliding). The high value of nCG for the unpressed aluminium (Figure 27) could represent a regime leading into a power-law breakdown (PLB) region at rapid strain rates and/or high stress levels.
\n\t\t\t\t\tThe analysis of Figure 27 indicates that creep in pure aluminium after ECAP occurs by the same mechanism as in conventional coarse-grained materials with intragranular dislocation glide and climb as the dominant rate-controlling flow process. Therefore, the activation energy for creep QC should be the same as the value of the activation enthalpy of lattice self-diffusion ∆HSD (∆HSD\n\t\t\t\t\t\t\n\t\t\t\t\t\t\t
It can be concluded that the creep resistance of high purity aluminium is increased considerably already after one ECAP pass. However, successive ECAP pressing leads to a noticeable decrease in the creep resistance. The results of microstructure investigations indicate that an inhomogeneity and thermal instability of the ECAP microstructure may strongly influence the creep behaviour of the pressed material [48].
\n\t\t\t\tThe combination of solid-solution strengthening and precipitation strengthening in creep of aluminium alloys at elevated temperatures has not been extensively studied. Numerous reports dealt with the creep behaviour of Al-Mg solid solution [68-70]. Most precipitation strengthened aluminium alloys currently being used are limited to relatively low temperature usage, because of the dissolution and/or rapid coarsening of their precipitates. An exception represents Al-Sc alloys containing low volume fractions of very fine coherent cuboidal Al3Sc precipitates. The dislocation creep behaviour of coarse-grained binary Al-Sc alloys at 573 K and the precipitation strengthening effect of the Al3Sc phase were investigated by Fuller et al. [71] and Seidman et al. [72]. Recently, the effect of Mg addition on the creep behaviour of an Al-Sc alloy was reported by Marquis et al. [31]. It was found that the creep strength of an Al-3wt.% Mg-0.2wt.%Sc alloy, containing Mg in solid solution and Al3Sc as nanosize precipitates, is significantly improved compared to binary Al-Sc alloys.
\n\t\t\t\tAs it was reported in Section 4.2. the processing by ECAP of a coarse-grained high purity aluminium provided a potential for marked improvement in the creep properties. Accordingly, Section 4.3. reports on a systematic study of the creep behaviour of the ECAP processed aluminium alloys containing low volume fraction of Al3Sc precipitates to elucidating the effect of ECAP on their creep resistance.
\n\t\t\t\t\t\n\t\t\t\t\t\tFigure 28a shows standard strain ε versus time t curves for the as-received (unpressed) Al-0.2wt.%Sc and Al-3wt.%Mg-0.2wt.%Sc alloys and those for the same alloys processed by ECAP through 8 passes at 473 K and 50 MPa (an exception is 80 MPa for Al-3wt.%Mg-0.2wt.%Sc alloy in the as-received state). The creep tests in compression were interrupted at a true strain of about ~ 0.35. These standard creep curves were replotted in the form of the instantaneous creep rate dε/dt versus time t as shown in Figure 28b. It is clear that no-one of the creep curves exhibits a well-defined steady state. In fact this stage is reduced to an inflection point of the dε/dt versus t curve. Supposing that the instantaneous creep rate dε/dt at given stress and temperature is a certain measure of the “softness” of the microstructure, then the dε/dt-t plots reveal the time evolution of this “softness”. However, the dε/dt-ε plots may give additional information, since they reflect the effect of the plastic creep strain on the instantaneous “softness” of the microstructure (Figure 28c).
\n\t\t\t\t\tCreep curves for specimens after ECAP processing through 8 passes and for unpressed specimens: (a) standard creep curve, (b) creep rate vs. time, (c) creep rate vs. strain.
The differences in the minimum creep rates for pure Al, Al-0.2wt.%Sc and Al-3wt.%Mg-0.2wt.%Sc in the as-received and as-pressed conditions are illustrated most readily in Figure 29 showing the variation of the minimum creep rate with applied stress. The results demonstrate that for pure aluminium at high stresses the minimum creep rate of ECAP material may be up to one order of magnitude lower than that of the unpressed material, although this difference decreases with decreasing applied stress so that, at 10 MPa is negligible. By contrast, when tests of Al-Sc and Al-Mg-Sc alloys are performed at the same stress, the creep rates in the as-pressed alloys are faster than in the unpressed alloys by more than two and/or three orders of magnitude on the strain rate scale. The stress dependence of the minimum creep rate for the as-pressed Al-Mg-Sc alloy at lower stresses (σ < 20 MPa) is different in trend, which is clearly demonstrated by the characteristic curvature on the plot in Figure 29.
\n\t\t\t\t\tStress dependences of minimum creep rate for pure aluminium and its alloys in the unpressed and ECAPed conditions.
The linear extrapolation procedure for determining the threshold stress.
Such behaviour is generally associated with the presence of a threshold stress marking a lower limit stress below which no measurable strain rate can be achieved [73,74]. The threshold stress value determined from linear plot [75] of \n\t\t\t\t\t\t\t
Recently the tensile creep experiments on the same binary Al-0.2%Sc alloy were reported [38]. The stress dependences of the minimum creep rates and the times to fracture for this alloy after 1 and 8 ECAP passes are shown in Figure 31. As demonstrates by Figure 31, the pressed alloy after 8 ECAP passes exhibits the very similar value of n as the results obtained by compression tests (Figure 29).
\n\t\t\t\tThe results from this investigation on precipitation-strengthened aluminium alloys do not confirm a general validity of the conclusion of our earlier results that processing by ECAP of a coarse-grained aluminium gave a potential for an improvement in the creep resistance [26,27]. By contrast, the Al-0.2wt.%Sc and Al-3wt.%Mg-0.2wt.%Sc alloys exhibited faster creep rate than their coarse-grained counterparts when creep tested under the same loading conditions.
\n\t\t\t\t\tStress dependence of (a) the minimum creep rates, and (b) times to fracture for Al-0.2Sc alloy.
The observed values of the stress exponents n = \n\t\t\t\t\t\t\t
The detailed analysis was undertaken by Kawasaki et al. [36] to examine the flow characteristics of the ultrafined-grained Al-3Mg-0.2Sc and Al-0.2Sc alloys [33-35,43]. The theoretical predictions of the minimum creep rates have shown that Nabarro-Herring creep is too slow to account for the creep deformation in an ultrafine-grained Al-3wt.%Mg-0.2wt.%Sc alloy [33] but there is good agreement, to within an order of magnitude, with the predictions of superplastic flow except only at the lower stresses where the points deviate from linearity and there is evidence for the presence of a threshold stress (see Figure 29).This threshold stress probably arises from the presence of coherent Al3Sc precipitates.
\n\t\t\t\t\tFor an Al-0.2wt.%Sc alloy crept in compression at 473 K [33] predictions for Nabarro-Herring diffusional creep were again too slow but there was excellent agreement between the experimental datum points and the predicted behaviour in superplastic flow. The slightly higher stress exponent after ECAP may reflect an inhibition in GBS at the lowest strain rates due to the presence of intergranular Al3Sc precipitates. Finally, the results on an Al-0.2wt.%Sc alloy tested in creep under tensile conditions at 473 K [43] again shown that the predicted behaviour for Nabarro-Herring creep is too slow but there is reasonable agreement with the model for superplastic flow.
\n\t\t\t\t\tThus, the results from both sets of experiments exhibit a general consistency with the predicted behaviour for conventional superplasticity. The theoretical predictions provide a clear demonstration that conventional creep mechanisms, already developed for coarse-grained materials, may be used to explain the flow characteristics of materials with ultrafine grain size. Furthemore, at least for the aluminium alloys examined in this Chapter, it is not necessary to involve any new and different creep deformation mechanisms.
\n\t\t\t\tCreep ductility is very important for various shaping and forming technological operations at elevated and high temperatures and especially for avoiding catastrophic failure in load-bearing parts of high temperature components. Creep strength and ductility are the key creep properties of creep-resistant materials but these properties typically have opposing characteristics. Thus, these materials may be strong or ductile but they are rarely both. In this connection, recent findings of high strength and good ductility in several bulk ultrafine-grained (UFG) metals produced by severe plastic deformation (SPD) are of special interest [4].
\n\t\t\t\tTypically, creep ductility in tension which can be characterized by the strain to fracture εf, is given by
\n\t\t\t\twhere σ is the applied stress, T is the absolute temperature, S is some parameter that characterizes the structure of specimen and tf is the time to fracture. According to the first mean-value theorem of a definite integral eq. (4) can be expressed as
\n\t\t\t\twhere \n\t\t\t\t\t\t
A strong influence of the microstructure on creep behaviour has been observed in various UFG materials [22,23,38,49,54,76]. By contrast, no report is available describing the link between microstructure and creep ductility in UFG materials processed by ECAP at elevated and high temperatures.
\n\t\t\t\tTo investigate a course of creep strain during creep exposure and for mutual comparison of the fracture strains Figures 32a and 33a can be replotted in the form of the instantaneous strain rate dε/dt versus creep exposure time t normalized to the time to fracture tf as shown in Figures 32b and 33b. The figures demonstrate that very significant strain contributions to the fracture strains are generated during the last tenth of creep life, however, an accumulation of the creep strains during the course of creep exposure slightly differ.
\n\t\t\t\tCreep curves of pure aluminium for unpressed state and various number of ECAP passes (creep in tension up to fracture): (a) standard creep curves, (b) creep strain ε vs. time t/tf.
A question naturally arises about an approach to the problem of creep ductility enhancement. In this connection different approaches may be considered [4]. It has been suggested that some ductility enhancement may be associated with an increase in the fraction of high-angle grain boundaries with increasing number ECAP passes and with a consequent change in the controlling deformation mechanisms due to the increasing tendency for the occurrence of grain boundary sliding. A possible explanation for the occurrence of intensive grain boundary sliding in UFG materials is that diffusion is more rapid in ECAP processed materials with highly non-equilibrium grain boundaries [77,78]. Accordingly, it appears that grain boundary sliding is easier in these UFG materials during creeping at ambient temperatures leading to the possibility of increased ductility. Figure 34a shows the measured strain to fracture as a function of the fraction of HAGBs when testing under creep loading conditions presented in Figures 32 and 33.
\n\t\t\t\tCreep curve of an Al-0.2wt.%Sc alloy for unpressed state and various number of ECAP passes (creep in tension up to fracture): (a) standard creep curves, (b) creep strain ε vs. time t/tf.
It is apparent from Figure 34a that specimens with higher fraction of high-angle grain boundaries exhibit higher ductility. Further approach to the problem of ductility enhancement at room temperature in ultrafine-grained materials was suggested through the introduction of a bimodal or multi-modal grain size distribution [79-81]. In this ductility strategy, the ultrafine-grained matrix in the bimodal microstructure provides the high strength, while the relatively large grains of the order of micrometers contribute to the ductility. The existence of large grains may also release stress concentrations, thereby delay the early fracture of the specimens, and allow further plastic deformation to take place in the ultrafine-grained matrix. Furthermore, the investigation of copper showed that bimodal structures may increase the ductility not only during tensile testing but also during cyclic deformation [82].
\n\t\t\t\tUnfortunately, the results of this work have not advocated the advantages of using material with a bimodal and/or multimodal grain size distribution for obtaining at the same time good creep strength and ductility at elevated or high temperatures. As it follows by inspection of Figures 19 and 20 typical feature of the boundary structure of UFG materials is its inhomogeneity; especially at low values of ECAP passes N completely different structures are observed. The dispersion of observed grain profile areas attains quite enormous values of the coefficient of variation CVa (between 6 to 16) in the whole range of N ≤ 12. In such circumstances, a great variability of times to fracture and creep fracture strain is a natural consequence of short as well as long-range inhomogeneity of microstructure of tested specimens. A very pronounced improvement of creep ductility and strong evidence for a significant role of microstructure homogeneity are documented in Figure 34b. These data indicate possible relationship between the ECAP microstructure and the creep ductility of the pressed materials [48]. The better is homogeneity of pressed material the better is its creep ductility. Thus, using a non-uniform grain structure by mixing ultrafine – sized and larger grains to derive enhanced ductility [83] seems not to be beneficial at elevated and/or high temperatures.
\n\t\t\t\tDependence of the strain to fracture εf on (a) the fraction of high-angle grain boundaries θ > 15°, and (b) on the coefficient CVa for EBSD range ∆ ≥ 15º. Crept at 473 K and 15 MPa (Al 99.99%) or 20 MPa (Al – 0.2% Sc) – see
Finally, due to a frequent use of miniaturized tensile specimens in research on UFG materials the specimen dimension and/or geometry effects could be considered for ductile behaviour of these materials. Thus, the thickness effect is mainly caused by the necking geometry and/or fracture modes, and the gauge length effect originates from the strain definition [84]. Therefore, no existence of a standardized protocol of creep specimen and a great variety of specimen sizes and geometries have been used by different authors, primarily depending on SPD techniques used and on the availability of material. Thus, the results creep ductile behaviour measured using such different specimens are hardly comparable.
\n\t\t\tHigh purity (4N) aluminium, a ternary Al-3wt.%Mg-0.2wt.%Sc and a binary Al-0.2wt.%Sc alloys were processed by equal-channel angular pressing (ECAP) through 1-12 passes and examined by TEM, SEM and EBSD microscopy. The microstructural investigations reveal that from 1 to 4 ECAP passes the microstructure evolves from elongated subgrains to an essentially equiaxed array of ultrafine grains (d < 1μm) and from 4 to 12 passes there is no substantial change in average grain size. The boundary misorientation angle and the fraction of high-angle grain boundaries (HAGBs) increase rapidly up to 4 ECAP passes and at a slower rate from 4 to 12 passes.
\n\t\t\tIt was found, that the creep resistance of pure aluminium is considerably increased in comparison with the coarse-grained unpressed state already after first ECAP pass. However, successive ECAP pressing lead to a noticeable decrease in the creep properties. This softening may be related to the decrease of the spacing of HAGBs at approximately constant subgrain size with increasing number of ECAP passes, resulting in the fraction of low-angle boundaries decreasing considerably. This indicates that HAGBs have lower strengthening effect under creep than low-angle ones. The softening by HAGBs may be explained in terms of the indirect effect which grain boundaries exert on the creep resistance by influencing the evolution of the dislocation microstructure in modifying the rates of generation and annihilation of dislocations. Further, a progressively increasing contribution of grain boundary sliding along HAGBs to overall creep strain can be expected as a consequence of accompanying transformation of low angle boundaries towards an equilibrium state when the number of passes increases. The results demonstrate that creep occurs in pure aluminium after processing by ECAP by the same mechanism as in conventional coarse-grained materials with intragranular dislocation glide and climb as the dominant rate-controlling deformation processes. Therefore, the higher creep resistance of pressed aluminium cannot be a consequence of any significant change in the rate controlling process.
\n\t\t\tThe present results concerning Al-3wt.%Mg-0.2wt.%Sc and Al-0.2wt.%Sc alloys do not confirm a general validity of the conclusion that processing by ECAP of a coarse-grained material gives a potential for an improvement in the creep resistance. By contrast, the minimum creep rates for the pressed alloys are more rapid than the rate attained in the coarse-grained state. The creep mechanism which most probably plays the dominant role in creep of the coarse-grained alloys is the dislocation climb-bypass mechanism in the presence of elastic interaction between dislocations and coherent Al3Sc precipitates. Strong support for making use of the same conventional creep mechanism in interpreting the creep characteristic of pressed alloy was presented. Although it seems acceptable to associate these faster creep rates with the smaller grain size after ECAP due to more intensive grain boundary sliding, there are probable further reasons strongly influencing an intragranular deformation mechanism which seems to be the rate-controlling process. First, the pressed alloys contain a high dislocation density because of the intense straining imposed during pressing. Second, an important strain contribution during creep of pressed alloys is produced by intensive mesoscale sliding of groups of grains along shear bands. Finally, the occurrence of nanopores formed during ECAP pressing at matrix/precipitate interfaces and resulting in the decohesion at coherent particle surface could have a strong effect on the creep behaviour of pressed alloys. It is important to note that the experimental results are in reasonable agreement with a theoretical model for superplastic flow that was developed earlier for coarse-grained alloys. This conclusion is important because of the many results now available demonstrating the occurrence of superplastic elongations in materials processed by ECAP and because of the new possibilities for achieving exceptional superplastic elongations in materials with ultrafine grain sizes.
\n\t\tThe authors acknowledge the financial support for this work by the Czech Science Foundation under the Grant No. P108/11/2260. This work was realized in CEITEC – Central European Institute of Technology with research infrastructure supported by the project CZ.1.05/1.1.00/02.0068 financed from European Regional Development Fund. We would like to cordially thank late Dr. Ivan Saxl of Mathematical Institute of the Academy of Sciences of the Czech Republic for his useful advices concerning the correct applications and interpretations of strereological methods for quantitative characterization of microstructure.
\n\t\t—Herophilus, 300 B.C., Greek physician and pioneer of anatomy.
\nWellness is an essential training for new physicians, not only for their own well-being but for the well-being of their patients and the communities to which they belong. The reason why this should matter is obvious: depressed residents make 6.2 times as many medication errors per resident month as residents who are not depressed [1]. As mentioned in the previous chapter, physician wellness is also a public health issue because many US patients lose their physicians to suicide every year [2]. Burnout and depression in physicians, residents, and medical students is something we can objectively measure, study, and intervene to prevent. Residency programs are gradually adopting the idea of wellness into their culture and curriculum, with some initial research beginning to show the positive effects. Although the objective is the avoidance of burnout, the focus truly needs to be on a cultural shift away from rigorous study that requires self-sacrifice and an extreme work-life balance mismatch to a culture that is focused on education and the support of a balanced workplace that promotes wellness. Tenets of wellness need to be woven into the fabric of medical education, just as clearly emphasized as those of professionalism and patient safety.
\nAs defined by Christina Maslach, burnout is a syndrome characterized by a loss of enthusiasm for work (emotional exhaustion), feelings of cynicism (depersonalization), and reduced personal accomplishment [3]. Although there has been much published on the subject of burnout in residency, there has been significantly less focused on the promotion of wellness [4]. Despite the fact that some programs have adopted wellness initiatives for their residents, many are still lacking a standardized, evidenced-based curriculum proven to improve outcomes. Similarly, some medical schools are beginning to integrate wellness practices into their curriculum, or at least into their extracurricular activities on offer. The issue of burnout must arise in medical school, as students entering medical schools are no different from their non-medical peers in terms of their mental health. Soon after starting their medical studies, the students’ overall mental health state declines [5]. Studies have shown by making simple modifications in medical schools such as transitioning to a pass-fail model, strengthening student relationships with faculty and peers, and incorporating mindfulness and stress reduction techniques into training, we can prevent burnout and help manage preexisting depression [5].
\nOne of the most intuitive places to start with a wellness curriculum is the encouragement and integration of physical activity into one’s routine. There are numerous studies in the general population on the psychological, emotional, and mental well-being benefits derived from physical activity and time spent in an outdoors environment. Some studies have investigated the prevalence and possible incentivization of exercise for those in residency. One study that looked at internal medicine residents showed that although 35% of residents exercised daily prior to starting residency, only 4% were found to exercise daily during residency. About 41% of residents failed to meet the recommended guidelines for exercise as defined by the Department of Health and Human Services (HHS), and 80% reported a decrease in physical activity since starting medical training. Of greatest interest is that those residents who met the HHS guidelines for physical activity were less like to be burned out than their fellow residents (OR 0.38) [6]. A recent meta-analysis of trials investigating the impact of exercise on depression demonstrated a standardized mean difference of −0.62 (a moderate effect) in depression at the end of the study when comparing exercise with no treatment or a control intervention [7].
\nA study at the Mayo Clinic in 2013 investigated the specific effect of exercise on burnout in post-graduate training. They invited residents and fellows to participate in a voluntary, team-based 12-week exercise program. Prior to the intervention, only 31% of medical trainees were meeting HHS recommendations for physical activity. At the completion of the program, rates of adequate physical activity among the participants significantly increased (to 48%) compared to non-participants (23%). Most importantly, there was an increase in quality of life and a decrease in rates of burnout [8]. Another study identified that there are specific benefits to be gained by exercising outdoors, particularly in an alpine environment. The study showed that this environment improved symptoms of depression and quality of life measurements in those suffering from depression. Those with major depression were found to have lower levels of resiliency than their health counterparts at baseline, and that being physically active in an alpine environment was positively associated with an increase in resilience [9].
\nThere has been significant research published to date on emotional intelligence (EI) and psychological resiliency as counterweights against workplace stressors and reactions to trauma in the workplace in nurses and other professions (not to mention the research on EI as a prognostic in various at-risk populations such as children with autism spectrum disorder). A study showed that there were individual positive curvilinear relationships between both EI and psychological resilience with post-traumatic growth in nursing students. Moderate-level EI and resilience were associated with the greatest growth [10, 11]. This echoes similar studies in caregivers, where it was found that lack of resiliency was one of the main predictors of burden [12]. A separate study that investigated EI, resilience and burnout in pediatric and medicine-pediatric residents further supported the above conclusions. They found that self-compassion and mindfulness were positively associated with resilience and inversely associated with burnout. Interestingly, physician empathy and EI were not significantly correlated with burnout or resilience, suggesting that the impact of mindfulness and resiliency on burnout may be the reasonable focus to emphasize during curriculum development [13].
\nFailed attempts in the past have revolved around the individualized approach. The institution may ask what is wrong with the individual physician that causes him/her to become burned out. Rather than focusing on
Similarly, building a community at work is exceedingly important to overall physician well-being. Creating the space for interaction through literal spaces like physician lounges or through protected time during work hours for interest group meetings can be of significant value. Shared professional identity and support from one’s peers have a demonstrated ability to reduce burnout [16]. An institution should be motivating its physicians to be efficient and provide good patient care, but sometimes these incentives can contribute to burnout and thus meet a self-defeating end. For example: the common payment structure of the base salary plus a productivity bonus. Productivity can be increased by seeing more patients per unit of time, ordering more tests, and by working longer hours. All three can contribute to reduced quality of care and increased physician burnout, despite being considered “productive” actions. The more recent shift in payment structure has been toward one of quality measures and patient satisfaction, but this model has yet to demonstrate actual improvement in quality and may not truly change the day-to-day experience of the overworked physician [17].
\nRegardless of the ultimate practices that they enter, we can significantly improve the outcomes if we integrate wellness into medical training as early as possible. There is a growing number of programs and health systems realizing this and creating wellness curricula for their residents. In the last few years, the American Medical Association (AMA) and the American Academic of Family Physicians (AAFP) have developed resources and even entire conferences devoted to the subject of physician wellness [18, 19]. More often than a few years ago, one sees the term “resident wellness” on the websites promoting residency programs. Not only is the concept of physician wellness becoming more common, but so is the development of interventions to try to increase wellness. One such study at the Mayo Clinic demonstrated that a simple intervention of biweekly 1-hour discussion groups focused on mindfulness, shared experience, and small-group learning not only improved engagement at work and decreased burnout, but the effect was sustained 1 year later [20].
\nA meta-analysis in 2016 evaluated over 2000 articles regarding physician burnout (not residents, specifically), with 15 randomized trials and 37 cohort studies. The literature indicated that multiple techniques could be utilized that resulted in clinically meaningful reductions in burnout. When examined grossly, these studies showed the rates of overall burnout, emotional exhaustion, and depersonalization decreased following intervention by 10, 14, and 4%, respectively [21].
\nThe following year, a meta-analysis was completed specifically to investigate the efficacy of interventions in reducing resident physician burnout. Approximately ½ of the studies included in the meta-analysis were specifically examining the effects of the 2003 & 2011 Accreditation Council for Graduate Medical Education (ACGME) duty hour regulations. The collective results showed that reduced work hours were positively associated with a decrease in emotional exhaustion and depersonalization and no effect on personal accomplishment. The same were also decreased with other interventions including self-care workshops and a meditation intervention [22].
\nThis collective research has been inspiring but less specific than desired to those in the medical education industry. Our aim was to implement a specific wellness curriculum, collect data illustrating its effectiveness, and present it to the readers in a comprehensive curriculum that is ready to implement. The data thus far collected has demonstrated the efficacy of our curriculum and we continue to strive to expand and add to the curriculum as the needs of our residents are established.
\nOur goal is to describe the most practical residency wellness education curriculum, to provide the tools for objective measurement of resident wellness and burnout, and to assist in career longevity of our young doctors. This should be a priority not only for those going through the process, but for patients and administrators, too. Physician burnout affects all groups and is a pervasive issue that affects healthcare at every level. Burnout leads to decreased quality of care, increased medical errors, decreased patient satisfaction, decreased productivity and time management, and a high rate of physician turnover [16]. It is in the best interests of patients, doctors, administrators and ultimately the larger community to devote time and resources to this issue.
\nWhenever an intervention or curriculum is devised, it is equally important to properly develop an assessment and a timeline of monitoring to see if the desired effect or educational goals are being achieved. The following assessments were found to be reliable, validated, and appropriate to use to evaluate burnout and work satisfaction in the physician [3, 23, 24].
\nThe original Maslach Burnout Inventory (MBI) is a self-assessment tool composed of 22 items pertaining to occupational burnout. This tool assesses the three categories of burnout: emotional exhaustion, depersonalization, and personal accomplishment. In our suggested curriculum, the abbreviated Maslach Burnout Inventory (aMBI) is used, which adds an additional category of assessment that indicates satisfaction with medicine. This inventory can be periodically administered to monitor for increasing or decreasing levels of burnout in the individual or group.
\nOf the 12 items in the aMBI, they are separated as follows:
\nEmotional exhaustion is measured by three items. Higher scores are correlated with feeling more emotionally overextended and exhausted at work.
\nDepersonalization is measured by three items. Higher scores are correlated with a more frequent sensation of detachment and unfeelingness toward the care one provides at work.
\nPersonal accomplishment is measured by three items. Lower scores are correlated with a reduced feeling of competence and success at work, and thus higher burnout.
\nSatisfaction in Medicine is measured by three items. Lower scores are correlated with a lower satisfaction in the career of medicine, and thus high burnout risk.
\nThe MBI and the aMBI have been shown to have strong reliability, as well as being highly applicable and valid specifically in evaluation of family practice physicians [3, 23, 25].
\nThe Utrecht Work Engagement Scale (UWES) was developed to measure the positive opposite of burnout, which is work engagement [18]. Work engagement is defined “as a positive, fulfilled, work-related state of mind that is characterized by vigor, dedication, and absorption.” The UWES was originally developed with 17 items, however multiple studies have since streamlined and verified the validity of the UWES-9 [24, 26, 27]. The nine-item UWES measures vigor, dedication and absorption with three items each as follows:
\nVigor: These three items aim to identify the subject’s vigor. A higher score indicates that the subject has a high level of energy and mental resilience while working and that they are willing to invest effort in their work. They also demonstrate persistence in the face of difficulty.
\nDedication: These three items aim to identify the subject’s dedication. A higher score indicates that the subject has a greater sense of significance, enthusiasm, inspiration, pride, and challenge in their work.
\nAbsorption: These three items aim to identify the subject’s absorption. A higher score indicates that the subject is more deeply engrossed in their work. While at work, they experience time passing quickly and may have difficulty in detaching themselves from their work [24, 26, 27, 28].
\nUWES is an important metric to use in conjunction with the aMBI as work engagement has significant associations with productivity and mental health. Work engagement has been associated with numerous benefits including mental and psychosomatic health, intrinsic motivation, and positive attitudes toward work and the organization and high performance [28]. Furthermore, engagement is a more chronic, persistent state that is not associated with the day-to-day variations in work activity as may be other measurable indicators [24, 26, 27, 28].
\nTransitioning from medical school training to residency can be very challenging. An important but often overlooked aspect of graduate medical education is to help residents develop strategies to improve and sustain their well-being so that they can not only successfully complete their training but enter practice ready to overcome challenges. This section presents successful, practical initiatives from our family medicine residency program to assess, monitor, and enhance resident physician well-being. We will walk you through the steps of our wellness curriculum development and implementation. We will also present the spreading change from a single family medicine residency program to the hospital health network.
\nIn August 2015, St. Luke’s Family Medicine Residency Program in Bethlehem decided to study its resident burnout and wellness in preparation for a faculty development on physician burnout. Residents were assessed with the abbreviated Maslach Burnout Inventory (aMBI) [23, 29, 30]. The aMBI assesses the following three components: emotional exhaustion, depersonalization, and personal accomplishment. The “satisfaction with medicine” component was added to see how all these risk factors affect the residents’ personal satisfaction with life in medicine.
\nOur analysis revealed that our residents’ burnout levels approximated the national average, with emotional exhaustion being the predominant burnout category. Once we had obtained baseline data, the next step was to develop an intervention for ongoing wellness reinforcement and burnout prevention. At a program retreat in September 2015, we introduced the topic of physician burnout and wellness. We also discussed ways to construct resiliency to combat burnout. We encouraged residents to make goals to improve their well-being in several dimensions including physical, emotional, spiritual, and support goals. At this retreat, we organized a working group composed of faculty and residents to address wellness as a way to combat burnout. The working group was led by a faculty member passionate about wellness and one volunteer from each class (six residents per class). Following the creation of this group, a wellness curriculum was developed to include: monthly wellness education, engagement opportunities, and recognition of resident/faculty accomplishments. The main objective was to obtain a baseline wellness/burnout level in the intern year followed by reassessment every 6 months through the entirety of residency training. Using this data, we aimed to shape the culture and wellness initiatives to the needs of our program.
\nBy the spring retreat in March 2016, we had made considerable progress in measuring wellness and burnout for our residents. Among other topics at the retreat, we presented “Practical wellness applications for our residency program to prevent burnout.” Given that the highest category of burnout was emotional exhaustion, we specifically targeted the topics chosen for our residency retreat and wellness curriculum to focus on decreasing this particular risk. Although we focused on emotional exhaustion, our retreat still addressed the other objectives of decreasing depersonalization, enhancing a personal sense of achievement, and implementing healthy habits to combat burnout.
\nThe wellness curriculum has two major components: hosting a residency retreat every 6 months and a monthly wellness meeting that was built into our previously established academic curriculum.
\nCommon topics addressed at the biannual retreat are as follows:
Communication and leadership skills
Conflict management exercise
Work-life balance and time management
Burnout and wellness in residency
Meaningfulness at work
The privilege of medicine
Strategies for day-to-day challenges
Teamwork and team building activities
Giving and receiving feedback
Every month, a 1-hour time block is reserved during academic day to have a focused discussion either in small groups (five–six residents each) or with the entire group (all residents). Topics presented focus on enhancing wellness and decreasing burnout focus on three main categories: (1) importance of physician wellness, satisfaction in medicine, and healthy lifestyle, (2) addressing emotional exhaustion, depersonalization, and (3) encourage sense of personal achievement. Individual session’s topics are summarized in Table 1.
\nImportance of physician wellness, satisfaction in medicine, and healthy lifestyle: \n
\n
\n
\n
| \n
Wellness curriculum topics: 1 hour presentation and discussion.
In the small group format, especially in topics like wellness goal and goal setting, we use coaching technique to skill set in prioritization, self-care, and self-awareness [31].
\nAlong with developing a curriculum to enhance wellness and decrease the risk of burnout, we also developed both a formal Physician Well-being Policy and a Resource Manual in March 2017 and May 2017, respectively. This was given to every intern at the beginning of their academic year to provide them the tools and resources they would need to optimize their wellness. In August 2017, we also added the Work Engagement Scale (UWES) to be used in conjunction with the aMBI [26].
\nSince August 2015, we have assessed the wellness/burnout of our residents every 6 months. Results were compiled looking at both our program as a group, as well as separated out by class. Of significant interest is that we were able to follow the class of 2018 from their intern year to graduation, thus collecting a full 3 years of data from the same cohort. Figures 1 and 2 and Tables 2 and 3 summarize the data for both the residency as a whole and for the class of 2018 from summer 2015 to winter 2018.
\nPercent at risk per burnout category over time: program.
Percent at risk per burnout category over time: class of 2018.
Program data | \nEmotional exhaustion | \nDepersonalization | \nLow achievement | \nLow satisfaction with medicine | \n
---|---|---|---|---|
Summer 2015 | \n41% | \n33% | \n4% | \n22% | \n
Winter 2015–2016 | \n56% | \n31% | \n12% | \n25% | \n
Summer 2016 | \n41% | \n18% | \n18% | \n18% | \n
Winter 2016–2017 | \n56% | \n50% | \n13% | \n19% | \n
Summer 2017 | \n44% | \n27% | \n27% | \n0% | \n
Winter 2017–2018 | \n39% | \n11% | \n12% | \n19% | \n
St. Luke’s family medicine residency program data.
Class of 2018 | \nEmotional exhaustion | \nDepersonalization | \nLow achievement | \nLow satisfaction with medicine | \n
---|---|---|---|---|
Summer 2015 | \n20% | \n20% | \n20% | \n40% | \n
Winter 2015–2016 | \n56% | \n31% | \n12% | \n25% | \n
Summer 2016 | \n40% | \n20% | \n0% | \n20% | \n
Winter 2016–2017 | \n50% | \n33% | \n0% | \n17% | \n
Summer 2017 | \n66% | \n16% | \n0% | \n16% | \n
Winter 2017–2018 | \n50% | \n0% | \n0% | \n0% | \n
St. Luke’s family medicine residency class of 2018 data.
When examining the data from all residents as seen in Figure 1, we see decreasing indicators of burnout in three of four categories: emotional exhaustion decreased from 41 to 39%, depersonalization decreased from 33 to 11%, low satisfaction with medicine decreased from 22 to 19%, and low achievement increased from 4 to 12%. For the most part, there is a seasonal bimodal effect exhibited in most of these categories with increased burnout occurring in the winter with subsequent improvement in the summer months.
\nIn Figure 2, the data from the class of 2018 was analyzed from the beginning of their intern year in 2015 until the winter before graduation in 2018. The risk factors of depersonalization, low achievement, and low satisfaction with medicine all decreased from 20–40 to 0%, while the emotional exhaustion category increased from 20 to 50%. Again, the seasonal bimodal pattern is noted in some risk factors, with peaks in winter and troughs in summer. Of note, during the residency training of this particular class significant stressors occurred including the departure of three residents during a 1-year time period as well as many interpersonal conflicts among residents.
\nThe Work and Well-being survey was added as an additional study parameter in the summer of 2017 using the Utrecht Work Engagement Scale (UWES) [26]. So far, we only have two data points for summer 2017 and winter 2017–2018. The highest risk group which is not engaged at work and at risk for burnout has been constant at 33%.
\nOur residency has been using validated tools to assess physician burnout and engagement by using the aMBI and UWES. Our wellness program is effective in addressing burnout and promoting wellness given the encouraging result in the residency data in decreasing physician burnout from 2015 to present time in 2018. This result helps to validate the effectiveness of our curriculum.
\nThe call for addressing wellness was stated in ACGME letter from August 2016 to the community: “While a commitment to resident well-being has been a shared focus for all of us engaged in graduate medical education, the topic of physician well-being has recently risen to the forefront within the broader medical community… Layered on top of these challenges is the stigma that many physicians in training and in practice encounter related to acknowledging that they need help, and to feeling supported and safe in asking for that help.” [32] In March 2017, the ACGME revised the new section VI in Common Program Requirement asking each residency program to address physician well-being [33].
\nThe implementation of a well-being program in residency does, however, come with challenges. It requires additional time and energy for the faculty wellness champion to rally support, train, organize, and execute the wellness program. Cost may also be an additional factor for the already tight residency budget.
\nIn our residency program, we were fortunate to have the endorsement of the residency Program Director as well as the Designated Institution Official (DIO) in piloting wellness initiatives at our family medicine residency. We were able to get the support of faculty to promote wellness and participate in the curriculum for the residents. There is 1 hour per month set aside for the residents to participate in the wellness curriculum, and the involved faculty members spent about 1 hour per month on planning and preparing for the month’s activities. Preparation for the biannual retreat (scheduled for a 5-hour block of time) usually begins 3 months prior to the retreat and requires an additional 1 hour per month toward preparations from the involved faculty. In the week leading up to the retreat, participating faculty may spend up to 3 more hours making specific preparations for organized discussions, team building games, other activities and food arrangements. Therefore, it takes about 40–50 hours per year for each faculty dedicated to wellness curriculum involvement.
\nIn terms of cost, our residency utilized low-cost resources and implemented creative plans for our wellness activities. Wellness activities such as a walk in the park or exercising in the office can be accomplished without cost. The speakers and group leaders for the wellness topics were individuals from our own faculty or physician experts outsourced from the hospital network, again eliminating cost. At each monthly wellness meeting, our residency provided healthful snacks such as fruits, vegetables, healthy dips, crackers, and water which cost about $25–$30 per month. The retreat is usually held at a local park pavilion which costs approximately $50 per retreat to rent. Food and materials necessary for each retreat costs approximately $300. Therefore, in an academic year with two residency retreats, 12 wellness monthly meeting, the average annual cost to our residency budget is about $1000 to support residency wellness.
\nGiven the success of our residency program in addressing resident physician wellness, the DIO and GME has formed a physician well-being subcommittee to share the well-being initiatives we have developed, and to help spread the positive change to other residency programs within the network. The success in enhancing physician wellness goes beyond the work of individual residency program. It requires a multilevel approach from individual residents starting at orientation with assessment and personal goal setting to enhance well-being, to the program recognizing, assessing, organizing, and addressing well-being for residents, to the network endorsement by GME leadership and expanding wellness initiatives throughout the health network.
\nOur residency program has seen many positive outcomes through promoting a culture of wellness and addressing physician burnout. Promoting physician wellness takes more than just a wellness curriculum. It requires efforts from individual physicians, residency faculty at the program level, and also support at institution level. On a national level, the ACGME and many other national organizations have called for awareness and continued support to promote physician well-being and decrease burnout risks. Ultimately, this will help to fulfill the quadruple aims of healthcare in enhancing patient experience, improving population health, reducing costs, and improving the work life of the health care providers.
\nAs we continue to strive for a better future, we hope to build upon the excellent response to the wellness curriculum in our residency over the past 3 years. We will need to continue improving and modifying our approach to best meet the need of our residents. The structure of our current curriculum, including both monthly sessions and the content of the retreat is drawn directly from similar ideas present in the current literature [34]. As we continue to grow, we may consider the addition of required reading about wellness or introducing apps for mindfulness to the residents [35]. In term of topics, we could add in modules as suggested by Arnold et al. such as physician suicide, “I need help”, or “dealing with difficult consultants and staff” [36].
\nAnother area of focus for development would be addressing resiliency. A meta-analysis identified that the factors conferring the largest increase in resiliency were higher self-esteem, most positive attributional style, and lower perfectionism. Other factors with moderate impact on resilience include lower trait reappraisal and higher EI. This study thus identifies avenues for further development of curriculum modules in order to bolster well-being during residency through resiliency training [37].
\nIn recent years, literature has focused on a discussion of work-life balance versus work-life integration [38, 39, 40]. It is important to realize that a well-balanced integration for professional and private life is an important goal for the new generation of doctor [41]. We can consider changing our curriculum topic of work-life balance to work-life integration to emphasize that the work life and personal life can mingle and integrate. The physician can find a sense of joy and enhance well-being when there are meaningful interconnections between the many existing roles in both work and personal life.
\nAmong the many factors that could help improving physician well-being, coaching has been an effective tool in improving general well-being of employees and reducing burnout [42, 43]. In our small group, we use coaching techniques to help residents with goal settings, self-awareness, and self-care. However, we can consider getting more formal training for faculty in coaching residents so that we can have similar training standards for our coaching techniques.
\nGiven the above future directions, we will need to ensure that we continue to assess our physician wellness and evaluating all the initiatives to make sure that our curriculum is effective and validated. By continuing to measure burnout and wellness with the aMBI and the UWES every 6 months, we can monitor the wellness of our residents as well as assessing the effectiveness of our curriculum in improving well-being through time. There are many other validated instruments to assess work-related dimensions of well-being such as Oldenburg Inventory, Physician Work-life Study’s Single-Item, Copenhagen Burnout Inventory, Stanford Professional Fulfillment Index, Well-Being Index, Patient Health Questionnaire-9 [44]. The choice of using any of these validated would be depending on the resource available for the residents or organization in term of cost and personnel to conduct the test and monitor the results.
\nFormal assessment of wellness and burnout and the development and maintenance of a wellness curriculum throughout medical education is essential to future physician well-being. Given that the major factors contributing to resident well-being are autonomy, competency, social connectedness and time away from work, we will need more than just the residency program to address these issues [45]. The promotion of resident wellness requires a multi-level approach involving the individual physician, the residency, and the systematic support at organizational and national levels to wholly address burnout and enhance well-being. The ACGME’s Clinical Learning Environment Review (CLER) program requires residency to monitor physician well-being and burnout. It is recommended that the CLER can go further to expect institutions to have initiatives to enhance resident wellness and increase engagement through identifying suboptimal aspects of the learning environment and enhance resiliency in residency [46]. The health of our physicians is a far-reaching issue and must be addressed. Their longevity, productivity and happiness are heavily impacted by wellness and rates of burnout and can be greatly shaped throughout their medical education journey. The foundation has been built, but we must continue to fervently pursue the establishment of wellness as a pillar of medical education.
\nWe would like to acknowledge the support of the St. Luke’s University Health Network, with particular attention to those in the St. Luke’s Family Medicine Residency Program and the Graduate Medical Education Department.
\nThe authors have no conflicts of interest to disclose.
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