Chemical composition and creep rupture strength at 600°C of the 9-12Cr heat-resistant steels from 1950 to 2005 [19].
\\n\\n
More than half of the publishers listed alongside IntechOpen (18 out of 30) are Social Science and Humanities publishers. IntechOpen is an exception to this as a leader in not only Open Access content but Open Access content across all scientific disciplines, including Physical Sciences, Engineering and Technology, Health Sciences, Life Science, and Social Sciences and Humanities.
\\n\\nOur breakdown of titles published demonstrates this with 47% PET, 31% HS, 18% LS, and 4% SSH books published.
\\n\\n“Even though ItechOpen has shown the potential of sci-tech books using an OA approach,” other publishers “have shown little interest in OA books.”
\\n\\nAdditionally, each book published by IntechOpen contains original content and research findings.
\\n\\nWe are honored to be among such prestigious publishers and we hope to continue to spearhead that growth in our quest to promote Open Access as a true pioneer in OA book publishing.
\\n\\n\\n\\n
\\n"}]',published:!0,mainMedia:null},components:[{type:"htmlEditorComponent",content:'
Simba Information has released its Open Access Book Publishing 2020 - 2024 report and has again identified IntechOpen as the world’s largest Open Access book publisher by title count.
\n\nSimba Information is a leading provider for market intelligence and forecasts in the media and publishing industry. The report, published every year, provides an overview and financial outlook for the global professional e-book publishing market.
\n\nIntechOpen, De Gruyter, and Frontiers are the largest OA book publishers by title count, with IntechOpen coming in at first place with 5,101 OA books published, a good 1,782 titles ahead of the nearest competitor.
\n\nSince the first Open Access Book Publishing report published in 2016, IntechOpen has held the top stop each year.
\n\n\n\nMore than half of the publishers listed alongside IntechOpen (18 out of 30) are Social Science and Humanities publishers. IntechOpen is an exception to this as a leader in not only Open Access content but Open Access content across all scientific disciplines, including Physical Sciences, Engineering and Technology, Health Sciences, Life Science, and Social Sciences and Humanities.
\n\nOur breakdown of titles published demonstrates this with 47% PET, 31% HS, 18% LS, and 4% SSH books published.
\n\n“Even though ItechOpen has shown the potential of sci-tech books using an OA approach,” other publishers “have shown little interest in OA books.”
\n\nAdditionally, each book published by IntechOpen contains original content and research findings.
\n\nWe are honored to be among such prestigious publishers and we hope to continue to spearhead that growth in our quest to promote Open Access as a true pioneer in OA book publishing.
\n\n\n\n
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Payan-Carreira",dateSubmitted:"April 21st 2020",dateReviewed:"September 10th 2020",datePrePublished:"October 8th 2020",datePublished:"January 20th 2021",book:{id:"8545",title:"Animal Reproduction in Veterinary Medicine",subtitle:null,fullTitle:"Animal Reproduction in Veterinary Medicine",slug:"animal-reproduction-in-veterinary-medicine",publishedDate:"January 20th 2021",bookSignature:"Faruk Aral, Rita Payan-Carreira and Miguel Quaresma",coverURL:"https://cdn.intechopen.com/books/images_new/8545.jpg",licenceType:"CC BY 3.0",editedByType:"Edited by",editors:[{id:"25600",title:"Prof.",name:"Faruk",middleName:null,surname:"Aral",slug:"faruk-aral",fullName:"Faruk Aral"}],productType:{id:"1",title:"Edited Volume",chapterContentType:"chapter",authoredCaption:"Edited by"}},authors:[{id:"38652",title:"Dr.",name:"Rita",middleName:null,surname:"Payan-Carreira",fullName:"Rita Payan-Carreira",slug:"rita-payan-carreira",email:"rtpayan@gmail.com",position:null,institution:{name:"University of Évora",institutionURL:null,country:{name:"Portugal"}}},{id:"309250",title:"Dr.",name:"Miguel",middleName:null,surname:"Quaresma",fullName:"Miguel Quaresma",slug:"miguel-quaresma",email:"miguelq@utad.pt",position:null,institution:{name:"University of Trás-os-Montes and Alto Douro",institutionURL:null,country:{name:"Portugal"}}}]},book:{id:"8545",title:"Animal Reproduction in Veterinary Medicine",subtitle:null,fullTitle:"Animal Reproduction in Veterinary Medicine",slug:"animal-reproduction-in-veterinary-medicine",publishedDate:"January 20th 2021",bookSignature:"Faruk Aral, Rita Payan-Carreira and Miguel Quaresma",coverURL:"https://cdn.intechopen.com/books/images_new/8545.jpg",licenceType:"CC BY 3.0",editedByType:"Edited by",editors:[{id:"25600",title:"Prof.",name:"Faruk",middleName:null,surname:"Aral",slug:"faruk-aral",fullName:"Faruk Aral"}],productType:{id:"1",title:"Edited Volume",chapterContentType:"chapter",authoredCaption:"Edited by"}}},ofsBook:{item:{type:"book",id:"1532",leadTitle:null,title:"Semiconductor Laser Diode",subtitle:"Technology and Applications",reviewType:"peer-reviewed",abstract:"This book represents a unique collection of the latest developments in the rapidly developing world of semiconductor laser diode technology and applications. An international group of distinguished contributors have covered particular aspects and the book includes optimization of semiconductor laser diode parameters for fascinating applications. \nThis collection of chapters will be of considerable interest to engineers, scientists, technologists and physicists working in research and development in the field of semiconductor laser diode, as well as to young researchers who are at the beginning of their career.",isbn:null,printIsbn:"978-953-51-0549-7",pdfIsbn:"978-953-51-4996-5",doi:"10.5772/1999",price:139,priceEur:155,priceUsd:179,slug:"semiconductor-laser-diode-technology-and-applications",numberOfPages:390,isOpenForSubmission:!1,hash:"67c029e3a582411c5f9ab3a7dc28884f",bookSignature:"Dnyaneshwar Shaligram Patil",publishedDate:"April 25th 2012",coverURL:"https://cdn.intechopen.com/books/images_new/1532.jpg",keywords:null,numberOfDownloads:58399,numberOfWosCitations:29,numberOfCrossrefCitations:5,numberOfDimensionsCitations:17,numberOfTotalCitations:51,isAvailableForWebshopOrdering:!0,dateEndFirstStepPublish:"May 2nd 2011",dateEndSecondStepPublish:"May 30th 2011",dateEndThirdStepPublish:"October 4th 2011",dateEndFourthStepPublish:"November 3rd 2011",dateEndFifthStepPublish:"March 2nd 2012",remainingDaysToSecondStep:"10 years",secondStepPassed:!0,currentStepOfPublishingProcess:5,editedByType:"Edited by",kuFlag:!1,biosketch:null,coeditorOneBiosketch:null,coeditorTwoBiosketch:null,coeditorThreeBiosketch:null,coeditorFourBiosketch:null,coeditorFiveBiosketch:null,editors:[{id:"106345",title:"Prof.",name:"Dnyaneshwar",middleName:"Shaligram",surname:"Patil",slug:"dnyaneshwar-patil",fullName:"Dnyaneshwar Patil",profilePictureURL:"https://mts.intechopen.com/storage/users/106345/images/2754_n.jpg",biography:"Dr. D. S. Patil has been graduated from Poona University with a rank. He received the M.Sc. degree in Electronics Science with a first class in 1986 from the Poona university department of Electronics-Science. He secured M.C.M. degree with A+ grade from Poona University and the Ph.D. degree in Electronics from the North Maharashtra University, Jalgaon [Maharashtra], India. He qualified state eligibility test in Electronics in 1995. Since 1991, he has been working in the North Maharashtra University, Jalgaon and presently working as a Professor. He secured high school scholarship, national merit scholarship and received Rashtriya gaurav award sponsored by India International Friendship Society. He successfully completed a major Young scientist project from Department of Science and Technology, India. His name has been considered in the Steering committee as a member for the International Conference on Nanoscience and Technology 2008, Colarado, United States of America, International vacuum Congress, China 2010. He worked on the various committees of the universities. He has published about 157 papers in reputed journals and proceedings of the conferences. His research interests include the computer simulation of semiconductor, nano and optoelectronics devices, nano-electronics, Materials development and characterization for the nano-technological and optoelectronics applications, process automation using advanced microcontrollers and embedded systems, organic electronics and computer simulation of nanostructures including quantum dots and superlattice. He has developed with his research student a novel model of probability density spreading in GaN quantum wells. He has developed with research students, computer controlled dip coating system and microcontroller based spin coating system for the deposition of nano-materials. He has guided many students for their innovative research. He visited France and Germany to attend international conferences and present his papers. Moreover, he visited Technical University, Zurich, Switzerland to know the various activities and research carried out in Electronics Technology department. He worked as a reviewer for many reputed international journals. He has delivered many invited talks and popular lectures. He developed the Electronics Practical laboratory and curriculum as a first member of Electronics Department and framed syllabus of M.Phil. (Electronics) and M.Sc.(Electronics). Despite of this, he taught various courses to M.Tech. (VLSI Technology), M.C.A and B.Tech.(Chemical Technology). Recently, his name has been considered in Marscue Who’s who in the world.",institutionString:null,position:null,outsideEditionCount:0,totalCites:0,totalAuthoredChapters:"0",totalChapterViews:"0",totalEditedBooks:"1",institution:{name:"North Maharashtra University",institutionURL:null,country:{name:"India"}}}],coeditorOne:null,coeditorTwo:null,coeditorThree:null,coeditorFour:null,coeditorFive:null,topics:[{id:"1226",title:"Optoelectronics",slug:"optics-and-lasers-optoelectronics"}],chapters:[{id:"35899",title:"Effect of Cavity Length and Operating Parameters on the Optical Performance of Al0.08In0.08Ga0.84N/ Al0.1In0.01Ga0.89N MQW Laser Diodes",slug:"effect-of-cavity-length-and-operating-parameters-on-the-optical-performance-of-al0-08in0-08ga0-84n-a",totalDownloads:3587,totalCrossrefCites:0,authors:[{id:"104427",title:"Dr.",name:"Alaa J.",surname:"Ghazai",slug:"alaa-j.-ghazai",fullName:"Alaa J. 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There is a worldwide need for the sustainability of current energy sources in order to ensure the viability of future generations, as long as these sources are environmentally friendly. In this sense, power plant designs for the future should ensure a cost-efficient reduction of CO2 emissions and improvements in efficiency of fuel consumption.
The essential function of a power station is to convert energy from fuel (fossil or nuclear) into electrical energy. In the steam power plant, this conversion involves consuming the fuel to produce heat which is then used to produce steam to drive a turbine. The mechanical energy of the turbine is then converted to electrical energy by an alternator. The steam temperature on the entrance of the turbine is essential to increase the efficiency of the conventional steam cycle.
The maximum steam temperature and pressure are limited by the performance of certain components. The main components which are critical are steam headers, superheater and reheater tubing in boilers, turbine valve chest, rotors and casings, main steam and reheat pipework, generator rotors, and bolts used for high-temperature applications. The boiler components are limited by corrosion and creep. Pipework also suffers creep as well as weld cracking and thermal fatigue. Turbine components are subjected to creep and fatigue (both thermal and mechanical).
Therefore, the development of improved structural materials to increase in thermal efficiency has been the driving force to develop new generations of 9-12Cr ferritic/martensitic (FM) steels [1, 2, 3]. The most relevant in-use properties that heat-resistant steels employed to manufacture components in power plants should fulfill are good mechanical properties, fabricability, corrosion resistance, and creep strength. As indicated above, creep strength has been the most studied and has led to innumerable research activities, aiming at improving the creep strength in 9-12Cr FM steel developments [4, 5, 6]. The disadvantage of these steels is their loss of strength beyond 600°C, so they need to be optimized to guarantee their use in the future power plants. In this chapter one of the most promising ideas described is applying a thermomechanical treatment (TMT) instead of a conventional treatment. The main contribution of the TMT is the ausforming, which, as other authors have reported, allows increasing considerably the number density of the thermally stable precipitates, i.e., MX nanoprecipitates. Consequently, the creep strength has improved greatly.
As it has been reviewed by Klueh in his seminal work on high-chromium FM steels [7], the design and production of 9-12Cr FM steels began in 1912 when Krupp and Mannesmann produced a 12 wt. % Cr steel containing 2–5 wt. % Mo. This type of steel was used for steam turbine blades, and it is still in use under the designation of X22CrMoV12. The 2¼Cr-1Mo bainitic steel grade normally known as ASTM Grade 221 (with nominal composition of Fe-2.25Cr-1.0 Mo-0.3Si-0.45Mn-0.12C) was firstly introduced in fossil fuel power plants in the 1940s and is nowadays widely used. The 9Cr-1Mo FM steel grade (known as Grade 9) is a natural evolution from Grade 22, seeking a better corrosion resistance and, hence, increasing the chromium addition. These two steel grades are the reference steels for heat-resistant application in power plants. Since then, the steady need of pushing up the operating conditions in conventional fossil-fired power-generation systems led to the development of several “generations” of steels with improved elevated-temperature strengths. The evolution of steel compositions (Figure 1), which began with G22 and G9 (zeroth generation) with 100,000 h creep rupture strengths at 600°C of about 40 MPa, has allowed for increased operating steam temperatures and pressures [1, 2, 3, 8, 9, 10, 11, 12]. Three generations of steels have been introduced since the introduction of G22 and G9, and a fourth generation is in development.
Flowchart showing the evolution of 9-12Cr FM steels [7].
The strategy adopted for improved corrosion and oxidation resistance for elevated-temperature operating conditions was the addition of carbide formers such as vanadium and niobium to add precipitate strengthening. Hence, the zeroth generation containing mainly 9-12Cr evolved to the 12Cr-MoV steels introduced in the power plants in the mid-1960s for thin- and thick-walled power station components. Their creep strength is based on solution hardening and on the precipitation of M23C6 carbides. These steels have been applied successfully in power stations over several decades [10]. These steels had increased 105 h rupture strengths at 600°C of up to 60 MPa (Table 1).
Chemical composition and creep rupture strength at 600°C of the 9-12Cr heat-resistant steels from 1950 to 2005 [19].
The second generation, developed in the late 1970s, is based on the modified 9Cr-1Mo, designated as G91 and HCM12 (see Table 1), which were developed for manufacturing of pipes and vessels for fast breeder reactors [10]. In this steel class, C, Nb, and V contents were optimized, N (0.03–0.05 wt. %) was added, and the maximum operating temperature increased to 593°C. The new steels have a duplex structure (tempered martensite and δ-ferrite). These steels have 105 h rupture strengths at 600°C of about 100 MPa. Of these latter steels, G91 has been used most extensively in the power-generation industry in all new power plants with operational temperatures up to 600°C [7]. The responsible mechanism for this substantial increment of creep strength as compared with 12Cr-MoV steels is the formation of thermally stable V and Nb carbonitrides. Besides lowering the Cr content down to 9 wt. %, tempered martensite microstructure also contributes to the higher creep strength [13, 14].
The Japanese steel development program led by Nippon Steel achieved the development of the P92 steel (NF616). This steel grade, designated as Grade 92, presents a further increases in stress rupture by the addition of 0.003 wt. % B and 1.8 wt. % W and reducing the Mo content from 1 to 0.5 wt. % [15, 16]. The addition of B ensures thermally stable M23(C,B)6 precipitates, whereas the higher W content leads to a higher amount of precipitated Laves phase [17, 18]. Grade 92, firstly introduced in the 1990s along with equivalent steel such as E911, fulfills the niche of steam operational temperature of 620°C for 104 h creep rupture strengths at 140 MPa.
Finally, the goal for the next steel generation being developed at present is pushing the limit of operation temperature above 650°C. This so-called fourth generation differs from the previous ones mainly by the addition of 3.0 wt. % Co as an austenite stabilizer because of the adverse effect of nickel on creep. They have projected 105 h creep rupture strengths at 600°C of 180 MPa [7]. In these steels with about 0.1 wt. % carbon, molybdenum has been further reduced or eliminated, and tungsten (2.6–3.0 wt. %) has been increased compared to third-generation compositions. In Table 1, an overview of the historical development of the 9-12Cr heat-resistant steels from 1950 to 2005 is shown.
Creep deformation is a thermally activated process, and the rate of deformation (creep rate) is extremely temperature sensitive. In metals, creep deformation becomes important at temperatures greater than about 0.4TM, where TM is the absolute melting temperature [20]. In the case of 9-12Cr FM steels, this temperature is approximately 450°C. Clearly, power plant materials operate in the temperature regime where creep process is significant. The creep properties of the material used limit the operating temperature of many power plant components, such as the turbines. Development of materials with an increased creep resistance is central to the use of power plants with higher steam temperatures.
Creep deformation can occur by a variety of different mechanisms. The mechanism that dominates depends on the stress and temperature conditions as well as the microstructure of the material.
In the case of power plant steels, the stress levels are relatively high, and the temperatures (compared with melting point) are relatively low. In the case of creep deformation, it is controlled primarily by dislocation movement and the thermal energy available for dislocations to overcome obstacles. A deformation mechanism map gives information about which mechanism will dominate for a particular set of conditions. Such a diagram for a G91 steel is shown in Figure 2. For the exposure conditions, for this material, a power law creep (dislocation creep) is expected to dominate.
Deformation mechanism map (D-MAP) of T91 steel calculated from experimental reported values [22, 23]. The red line indicates the experimental conditions considered in this work.
Power law creep involves the movement of dislocations, and the creep rate is a result of the balance between work hardening and recovery. Work hardening results in an increase in the dislocation density, while recovery leads to a reduction in the dislocation density. If the dislocation density remains constant, then the creep rate is given by Norton’s law [21]:
In this equation
A study of the possible creep mechanisms suggests microstructures would be expected to have good creep resistance under conditions used in power plant. In general, creep-resistant alloys are based on a matrix which is a solid solution. The presence of misfitting solute atoms in solid solution makes the passage of dislocations through the matrix more difficult. However, the majority of the creep resistance, at least in the early stages of service, is derived from precipitate particles. Ideally these particles should be small, and they should be widely and homogeneously distributed in large numbers through the matrix. The particles need to be stable at operating temperatures for which the alloy is designated, and they should be resistant to coarsening, as this will reduce their effectiveness as strengtheners. In general, excessive work hardening and very fine grain sizes, which provide strengthening at ambient temperatures, are considered detrimental in high-temperature alloys. This is because both of them provide easy diffusion paths and therefore lead to an increase in the creep rate.
The new environmental regulations and commercial needs of the industry are the driving force for the development of new heat-resistant steels that push forward the operational limits of current steels. In this framework, the high-Cr FM steels applied as structural materials in fossil-fired and in nuclear power plants need to implement the operating temperatures above 650°C [1, 2, 3, 8, 9, 10, 11, 12]. The mechanism responsible for creep strengthening in these steels is the solid-solution and dispersion strengthening.
In the particular case of the so-called 9Cr FM steels, the creep degradation is a consequence of the thermal evolution of their hierarchal martensitic microstructure constituted by prior austenite grains, martensitic packets, blocks, and laths [24]. The microstructural degradation during creep consists of the coarsening of the lath structure [12, 24]. Such coarsening is governed by the subgrain boundary formation and evolution inside the laths, which can be prevented at high temperatures, and virtually frozen, by the dispersion of proper precipitates. The precipitates pin boundary migration and dislocation motion, slowing down the degradation of the martensitic microstructure and hence reducing creep rates [25, 26].
There are two main actors for the microstructural stability driven by precipitation in 9Cr FM steels: The first one is the coarse M23C6 carbides located mainly at the grain boundaries either from the prior austenitic grains or from the blocks or martensite lath boundaries. The second one is the V- and Nb-rich MX carbonitrides. Contrary to M23C6, those MX precipitates are homogeneously distributed within martensite laths. Therefore, the ideal situation would consist of reducing the presence of the M23C6 carbides to the minimum since their fast coarsening induces crack formation at the particle-matrix interface and promoting the formation of MX carbonitrides (nanometric in size), since they will delay the lath coarsening as mentioned above; it has been studied extensively [27, 28, 29].
Thermomechanical processing of 9Cr FM steels has been revealed as a promising tool to promote a high number density of MX carbonitrides [30, 31, 32, 33, 34, 35, 36, 37, 38]. TMT involves different steps that need to be optimized to produce the most favorable precipitate microstructure for elevated-temperature strength.
The creep behavior of a material may be characterized by a number of different parameters which can be measured by performing the appropriate creep test. For metallic materials most creep tests are conducted in uniaxial tension with a dumbbell-shaped specimen similar to that used for tensile testing. The tests are carried out at a constant temperature and under either a constant load or stress. Applying a constant stress is more useful if the test is being employed to provide information about a creep mechanism.
The conventional treatments (AR) and TMT considered in this work were carried out on 10 mm in length and 5 mm in diameter cylindrical samples using a DIL 805A/D plastodilatometer (TA instruments) as described elsewhere [39, 40]. Due to the limited amount of material available after the TMT is carried out in the plastodilatometer, the creep properties were investigated by means of the small punch creep test (SPCT) performed at 700°C as it has been previously reported [41, 42]. The SPCT samples were cut transversally, from cylindrical specimens, with a thickness of 600 μm and a diameter of 8 mm. Then, the disks were ground on both sides down to a final thickness of 500 μm. In the setup of the SPCT, the lower and upper dice are connected via a thread to ensure the clamping of the sample. The load is applied by a ceramic punch ball which is in contact with the sample. A plunger rod is used to transmit the dead weight load to the punch ball. All these components are made of Al2O3 ceramics. The clamping device is surrounded by an electrical heater and a thermal insulation. The upper plate carrying the additional dead weight is guided by two pillars with ball bearings. The temperature is measured in the lower die directly under the sample. The displacement is measured by a capacitive sensor between the upper plate and the thermal insulation with an accuracy of ±1 μm. A load cell is placed between the upper plate and the plunger rod.
The disk deflection vs. time resulting from the SPCTs might be divided into three different regions similarly to conventional strain vs. time creep curves obtained from uniaxial testing. However, the failure in SPCTs occurs away from the load line with cracks propagating in a circumferential direction due to membrane stretching. Therefore, the first part of the disk deflection vs. time curve corresponds to the loading region where the spherical indenter loads on a very small contact area of the sample are. Since the stresses will be higher than the yield stress of the material, local plasticity and an initial large deformation are produced. This large deformation is accumulated in a short period of time. The second stage corresponds to the steady-state region, which coincides with most of the sample life, where the disk deflection rate reached almost a minimum. Finally, the third stage consists of an acceleration of disk deflection and fracture region. The interpretation of this behavior is that once a crack propagates to a critical length, the sample is no longer in balance, leading to an increase in deflection rate and to a reduction in the structure stiffness in the tertiary region. Another explanation might be due to the localized necking without crack presence. The deformation mechanism in the tertiary region is a mixture among accumulation of creep damage, geometric softening, and crack growth effect.
As it has been introduced in previous sections, the pioneer commercial 9-12Cr steels present an upper service temperature of 540°C, which was successfully increased in the late 1970s up to 595°C with the introduction of vanadium and niobium microalloying in the composition of the steel. This steel was used as a benchmark for the development of steels with upper-use temperatures of 600–620°C.
However, it is difficult keep pushing the higher operating temperature too much. Therefore, to continue to exploit the advantages of ferritic steels, oxide dispersion-strengthened (ODS) steels [43, 44, 45, 46] were introduced. The first successful alloy was presented in the 1960s, and, since then, it has been an active research field. ODS steels are strengthened by small oxide particles, but the complicated and expensive manufacturing route avoided the full implantation as structural material in the current power plants.
Despite being around for about 40 years, the ODS steels are still in the development stage because of having mechanical property anisotropy [43, 45, 47, 48, 49]. Therefore, an alternative strategy to achieve a high number density of precipitates is needed. In this section, we present preliminary results that allow us to conclude that conventional thermomechanical control processing strategy is adequate to achieve dispersion-strengthened steels.
Lath martensite is a particular microstructure that ensures microstructural stability. Furuhara and Miyamoto [50] described the variety of crystalline size in lath martensite structures. A hierarchy of lath martensite structure is clearly identified particularly in low-carbon steels. A prior austenite (γ) grain is divided into “packets,” each of which consists of a group of martensite laths with the same parallel close-packed plane relationship in the Kurdjumov-Sachs (K-S) orientation relationship, denoted as “CP group” recently. In general, a packet is partitioned into several blocks, each of which contains laths of a single variant of the K-S relationship. Blocks and packets are mostly surrounded by high-angle boundaries, whereas lath boundaries inside a block are of low-angle type.
The microstructure resulting from conventional industrial heat treatment consists of tempered martensite, which presents elongated subgrains with an average size of 0.25–0.5 μ m (Figure 3). Two types of precipitates, M23C6 carbides rich in chromium and MX carbonitrides rich in V or Nb, are present in the microstructure. The size of M23C6 carbides is around 100–200 nm, and they are precipitated on subgrain boundaries and prior austenitic grain boundaries. The size of MX carbonitrides is much smaller than M23C6 carbides, 20–50 nm, and they are in the matrix [51]. The purpose of this work is to produce a dispersion of nanosized precipitates by a controlled TMT, bearing in mind that a high number density of fine MX precipitates (Nb-MX and V-MX) should display superior high-temperature performance.
(a) Resulting hierarchy microstructure achieved by conventional heat treatment; (b) and (c) SEM micrographs of the as-received state; (d) and (e) TEM micrographs. Arrow heads point out the location of the M23C6 carbides on lath boundaries and MX carbonitrides within the laths [37].
The effect of austenitization temperature on the temper microstructure of G91 steel is analyzed in this section. Figure 4 schematically illustrates the two alternative processing routes considered:
High austenitization temperature (HAT): In order to achieve an almost complete solid solution in austenite of most of the potential MX precipitate formers, the austenitization condition set will imply an elevated temperature.
Thermomechanical treatment: The combined effect of the elevated austenitization temperature and a subsequent deformation will be studied with the aim of optimizing the MX-nanoprecipitate distribution during tempering of the martensitic microstructure.
Thermomechanical treatments investigated in this study [40].
For the sake of comparison, Figure 4 also includes the industrial manufacturing conditions for G91 steel named as-received (AR) condition. The goal of exploring the effect of austenitization temperature on the microstructure is to enhance the precipitation of nanoparticles during the subsequent tempering stage indicated in Figure 4. As it was mentioned above, the main cause for creep softening in conventional G91 is due to the recovery of the martensitic lath microstructure because of mechanisms, such as the dislocation movement, controlled by diffusion [12, 27]. The dislocation pinning by nanosized MX precipitates can delay this phenomenon, since they present an enhanced ripening resistance [8, 52, 53, 54]. The goal of undergoing such elevated temperatures in the HAT treatment as compared to conventional austenitization heat treatment (AR treatment) is to dissolve all the primary carbides in the microstructure and drive to solid solution all the potential carbide former elements. Therefore, the martensite formed after quenching from such elevated austenitization temperature keeps in solid solution most of the precursor elements of MX carbides (M = Nb, V; X = C, N) that might precipitate during the subsequent tempering.
It is important to consider that the austenitization temperature has to be high enough to eliminate as much as possible the primary carbides formed during the casting process, but lower than the delta ferrite formation temperature, in order to avoid the detrimental effect of this phase from a long-term creep property point of view. Computational thermodynamic calculations by means of Thermocalc® determine the optimum austenitization temperature in 1225°C (Figure 5).
Temperature evolution of phase mole fraction in G91 calculated by Thermocalc® [38].
The interest of TMT relies on the role that austenite deformation has on refining the martensitic microstructure [55, 56]. Depending on the deformation temperature, several are the mechanisms that affect the austenite microstructure, and hence, that could be transferred to the martensite upon quenching. If deformation temperature is above the non-recrystallization temperature, the freshly formed austenite microstructure will present a significantly reduced grain size that would induce the concomitant martensitic microstructural refinement. Similarly, by applying plastic deformation to the austenite at temperatures below the non-recrystallization temperature, which is the so-called ausforming processing [57], an austenitic microstructure with a high population of deformation bands will be formed. This would directly induce the preferential formation of some specific martensitic variants upon austenite transformation (martensite variant selection), leading to the development of strong transformation texture.
Figure 6 illustrates the IPF maps, SEM and TEM micrographs after HAT and TMT processing routes, and the reference (AR) condition. The first conclusion obtained is the coarsening of the block size (white arrows in Figure 6) in HAT and TMT conditions as compared with AR condition, because of the high austenitization temperature. Block widths of 2.7 ± 0.2 μm for AR condition were obtained; meanwhile, values of 4.12 ± 0.37 μm for HAT and 3.21 ± 0.27 μm for TMT were measured. The coarser the prior austenite grain, the coarser the block size. However, it is worth noting that finer block size is observed after TMT than with HAT, which is consistent with the fact that thermomechanical processing increases the low-angle substructure and decreases the block size of as-quenched martensite.
Martensite matrix, M23C6 precipitate, and MX-nanoprecipitate distributions after the different thermomechanical and heat [40].
The dislocation density after HAT and TMT was measured by XRD [40]. The results show a dislocation density of (14 ± 0.1) × 1014 m−2 and (28 ± 0.1) × 1014 m−2 after austenitization and ausforming, respectively. One might conclude from these results that the dislocation density in the as-quenched martensite after the TMT is substantially increased as compared with conventional treatment. A similar effect of the ausforming on the dislocation density was reported by other authors [58, 59]. Finally, TEM examination of the microstructure allowed us to determine the lath width of the martensitic microstructure. Values of 360 ± 35 nm for AR condition, 350 ± 20 nm for HAT condition, and 318 ± 32 nm for TMT condition were obtained, which are significantly finer than those reported after conventional treatments, i.e., lath size ranging from 300 to 500 nm [60].
The distribution of M23C6 precipitates in the tempered martensitic microstructure is also worth analyzing. Figure 6 illustrates the distribution of M23C6 carbides after AR, HAT, and TMT processing routes. Coarse and closely spaced M23C6 carbides, about 70 to 500 nm, were observed. The number density and average particle size of these carbides were determined by studying several SEM micrographs to determine values of 6.19 × 1019 m−3 and 141 ± 3 nm for AR condition, 8.24 × 1019 m−3and 124 ± 3 nm for HAT condition, and 4.11 × 1019 m−3 and 143 ± 5 nm for TMT steel. These values are very similar to those reported by Klueh et al. for the steel after conventional heat treatment [35].
On the contrary, the finely dispersed MX nanoprecipitates present inside the martensitic laths and associated with dislocations are also observed in Figure 6. Therefore, this result suggests the role of dislocations as potential nucleation sites for MX nanoprecipitates. Hence, the importance of ausforming in generating a homogeneous distribution of nanosized MX particles in the microstructure might be also foreseen. These spherical MX nanoprecipitates had a mean particle size of 12 ± 1 nm with a number density of 7.20 × 1021 m−3 for HAT steel and 9 ± 1 nm with a number density of 1.86 × 1022 m−3 for TMT steel. The MX precipitates are, in both cases, significantly smaller than those measured after AR condition, i.e., particle size of 25 ± 5 nm with a number density of 8.14 × 1019 m−3. The size values obtained after HAT and TMT are smaller, and the number density higher, than measurements reported in the literature after conventional heat treatments [61].
Figure 7 shows the disk deflection versus time curves obtained for the three conditions studied (AR, HAT, and TMT) at 700°C with a load of 200 N. The curves exhibit the three stages of creep that were described in previous sections. The first stage corresponds to the loading region where the spherical indenter loads the sample, and the mode of deformation is by bending. The second stage is characterized by a decrease in deflection rate and corresponds to the steady-state region with a minimum disk deflection rate. Finally, the third stage consists of an acceleration of disk deflection and fracture region. In the secondary and tertiary stages, stretching is the prominent deformation mode. Once a crack propagates to a critical length, the sample is no longer in balance, leading to an increase in deflection rate and to a reduction in the structure stiffness in the tertiary region until the final fracture.
(a) SPCT curves measured for the samples after the different thermomechanical and heat treatments and the creep fracture micrographs for the (a) AR, (b) HAT, and (c) TMT [40].
As indicated above, the minimum disk deflection rate (δd ) is an important parameter that can be evaluated by SPCT. The evolution of disk deflection rate with the applied load might be described by an equivalent expression to the conventional Norton’s power law for creep, which is similar to the expression used in Eq. (1):
where A is a temperature-dependent constant, F is the force applied on the specimen, and n is the force exponent. One might conclude, therefore, from Figure 7 that the creep strength has significantly improved after the TMT condition. The time to rupture was 2.5 and 1.24 times greater than AR condition, from 38 to 95 h and 48 h for the TMT and HAT, respectively. The δd was 2.9 μm·h−1 for the TMT sample, while for the HAT sample, it was 3.7 μm·h−1. These minimum disk deflection rates were significantly slower than the minimum disk deflection rate measured for the G91 in the AR condition, which was 9.5 μm·h−1.
The results obtained suggest that the increase in the number density of MX precipitates enhances the strengthening capability at high temperature, since they are able to pin more effectively the dislocations. Hence minimum creep rate is reduced and the onset of tertiary creep is retarded. The differences in minimum disk deflection rate and time to rupture between the sample after TMT and HAT support the importance of ausforming on improving creep resistance.
The next stage in the TMT after austenitization is the ausforming as shown in Figure 4. The effect of ausforming on low-carbon lath martensitic microstructure has been already described by Miyamoto et al. [62]. The authors reported that martensite variants with habit planes that are nearly parallel to the close-packed primary and secondary slip planes in austenite transform preferentially, i.e., martensite habit planes such as (575)γ that are nearly parallel to (111)γ and (−111)γ in asutenite [63]. Since strain is accumulated preferentially in (111)γ and (−111)γ slip planes during ausforming, this results in an increasing number of dislocation that might be transferred to martensite (011)M planes. Therefore, ausforming might increase the dislocation density in the resulting martensitic microstructure.
On the other hand, Takahashi et al. [64] reported recently the formation of Nb-cottrell atmospheres in low-carbon Nb-microalloyed steels. The authors explained that this mechanism is based on the fact that segregation energy of Nb to edge dislocation core was almost the same as the energy for grain boundary segregation. Besides, the large attractive interaction between Nb and dislocation core was due to its large atomic size. Therefore, such interaction between Nb atoms and dislocations retards the recovery of dislocation at high temperatures and, hence, stabilizes the microstructure at high temperatures. It might be expected that Nb presents the same behavior in the studied steel, preventing recovery after ausforming and promoting the fine and homogeneous MX carbonitride precipitation during tempering accordingly.
In this work the role of ausforming temperature by selecting 600 and 900°C, at a constant deformation of 20% (Figure 4), is explored. As mentioned above, the dislocation densities were estimated by XRD in fresh martensite after each ausforming condition studied [41]. Values of (2.8 ± 0.1) × 1015 m−2 and (1.9 ± 0.1) × 1015 m−2 were obtained for the ausforming at 600 and 900°C, respectively. These results show that the lower the ausforming temperature, the higher the dislocation density introduced in austenite is, which might be due to the fact that some of the dislocations in fresh martensite are inherited from deformed austenite as it was mentioned above.
On the other hand, Bhadeshia and Takahashi reported [65] an expression that allows to estimate the dislocation density (ρd):
This expression is valid only when the martensite start temperature (T) is between the range 297 and 647°C.
Extracting the data of the martensite start temperature from a previous work [38], the estimation of the dislocation density obtained after the different ausforming conditions can be estimated. In this sense, ausformed samples at 600°C present a martensite start temperature of 338°C; introducing this value in Eq. (3), a dislocation density of 5.97 × 1015 m−2 is calculated. Similarly, for the material ausformed at 900°C with a martensite start temperature of 374°C, the dislocation density calculated is 4.62 × 1015 m−2. These results are in the same order of magnitude than those measured by X-ray diffraction, which demonstrate that the ausforming increases the dislocation density in the martensite.
During the final stage (tempering), MX carbonitrides and M23C6 carbides precipitate, and the recovery of dislocations takes place. Because of the higher dislocation density of ausformed samples, the number density of finer MX increases, and these precipitates are found homogeneously distributed within laths, as it can be seen in Figure 8(a) pointed out by white arrows.
(a) MX carbonitrides (white arrows) within laths after thermomechanical treatment ausformed (20%) at 900°C; (b) size distribution of MX precipitates in the TMT samples for the two ausforming temperatures: 600 and 900°C [41].
The number density of MX precipitates (N) was determined through the direct measurements of spacing (λ) between MX carbonitrides from several TEM micrographs as indicated by Eq. (4):
Figure 8(b) shows the size distribution of the precipitates in ausformed material. In the material ausformed at 600 and 900°C, the average size of MX carbonitrides was 5.6 nm and 7.4 nm, respectively. The number density of MX carbonitrides was 9.39 × 1022 m−3 for the material ausformed at 600°C and 6.4 × 1022 m−3 for the material ausformed at 900°C. On the other hand, the reported values of the size and number density of MX carbonitrides after the conventional processing were 30 nm and 1020 m−3, respectively [61]. It might be concluded that ausforming promotes a refining of precipitates, up to five times as compared with conventional processing, as well as an increase in number density up to two orders of magnitude. In fact, these number densities and precipitate sizes are very similar compared to those corresponding to oxides present in oxide dispersion-strengthened (ODS) steels [66, 67].
The elevated number density of nanosized MX precipitates has a direct impact on creep response of this material as it can be clearly observed in Figure 9. This figure shows characteristic SPCT curves at 200 N, exhibiting the variation of specimen deflection with time. It might be concluded from this figure that introducing an ausforming step improves the δd significantly, and most precisely, the lower the ausforming temperature, the lower the δd is, and, hence, the better the creep resistance is.
SPCT curves for all samples tested at 700°C with a load of 200 N [39].
Scanning electron microscopy (SEM) images of fractured SPCT specimens for different conditions are shown in Figure 10. Radial cracks can be observed in all the TMT samples (Figure 10a–c). This is an evidence of the loss of ductility and indicates a brittle fracture, which is a change in rupture ductility in comparison to the conventionally treated sample. Those samples do not show radial cracks (Figure 10d). Besides, a higher reduction in thickness is evident in the conventionally treated sample in comparison to the TMT ones, suggesting a ductile fracture behavior.
Scanning electron microscopy images of the SPCT fracture surfaces for samples tested at 700°C with a load of 200 N: (a) G91-TMT 900_20; (b) G91-TMT 600_20; (c) G91-TMT 900_40; and (d) G91-AR [39].
To clarify the failure mechanisms, the fractured samples were cut and prepared adequately. Figure 11(a) and (b) shows the SEM images for the TMT samples ausformed at 600°C with a deformation of 20% and ausformed at 900°C with a deformation of 40%. It is worth noting in those images the existence of cavities nearby coarse particles, which are located at the vicinity of PAGBs. The EDS spectrum shown in Figure 11(c) allows us to conclude that these particles are M23C6 carbides with M = (Fe, Cr, Mo).
Scanning electron microscopy images: (a) M23C6 precipitates located at a prior austenite grain boundary in sample G91-TMT 600_20. The prior austenite grain boundary in this image has been highlighted with a dash line as a guide to the eye; (b) cavities associated with coarse M23C6 precipitates have nucleated at a prior austenite grain boundary in sample G91-TMT 900_40. Cavities have been pinpointed with arrows and (c) EDS analysis of the particle marked with a red arrow in image (a), close to a cavity [39].
The greater size of the M23C6 carbides at the vicinity of PAGB contributes to the inhomogeneous and localized deformation experienced by the TMT samples at these locations during creep. The local creep concentration close to PAGB would be promoting the nucleation of cavities that lead to the intergranular fracture with the brittle behavior.
Figure 12 shows different inverse pole figure (IPF) maps for all the samples under study before and after SPCT. It should be pointed out that, contrary to the lath boundaries that are not correctly indexed due to the step size used for the EBSD mapping, the block boundaries before and after SPCT are clearly disclosed. It is observed that the microstructures of the samples exhibit the characteristic lath-like morphology of the martensitic microstructure. However, such morphology is blurred in samples ausformed at 600°C because of the high deformation accumulated in the austenite during ausforming [38]. After the SPCT, it is observed that the original lath-like morphology has partially disappeared, and it has evolved towards a fine-grained equiaxed ferritic matrix. One might conclude from the microstructural observations made after SPCT that newly formed equiaxed grains are distributed homogeneously in the conventionally treated sample (AR), while these grains are located mainly nearby the prior austenite grain boundaries in the TMT samples, which is consistent with the fact that it is in these samples where the deformation accumulated is larger during creep.
Representative inverse pole figure (IPF) maps of the initial and after SPCT microstructures: (a) G91-AR; (b) G91-TMT 900_20; (c) G91-TMT 600_20; and (d) G91-TMT 900_40. Cavities are in white in post-SPCT microstructures [39].
Therefore, taking into account the results shown previously in the SEM micrographs (Figure 11(a) and (b)), the microstructural degradation would be a combined consequence of the accumulation of dislocations at the low-angle boundaries and the stress concentration close to the coarse M23C6 carbides, which lead to the progressive loss of the lath-like martensitic microstructure, which evolves to an equiaxed ferritic matrix. As it has been discussed above in the case of the TMT samples, the nucleation of cavities takes place close to M23C6 precipitates located at the prior austenite grain boundaries. The coalescence of the cavities formed surrounding the M23C6 carbides would initiate the cracks, and they will propagate along the prior austenite grain boundaries.
The more homogeneous distribution of the M23C6 precipitates in the conventionally treated sample favors the apparition of equiaxed grains in the whole martensitic matrix and develops the nucleation of cavities intragranularly, which provokes the transgranular fracture. Besides, in the TMT samples, the high austenitization temperature produces an enormous prior austenite grain sizes with concomitant large grain boundary surfaces, facilitating an earlier formation of the critical crack length that causes the brittle fracture [68, 69].
Effect of austenitization temperature: compared to the conventional heat treatments, the use of a higher austenitization temperature (1225°C rather than 1040°C), combined with an ausforming processing step at 900°C, allows the increase of the number density of MX precipitates up to three orders of magnitude after the tempering step, which raises the strengthening capability of the MX at 700°C up to 6.5 times. These microstructures have reduced considerably the minimum disk deflection rate and showed greater time to rupture during the SPCT carried out at 700°C. By contrast, such elevated austenitization temperature induces an important drop in ductility.
Effect of ausforming: the SPCT was applied to evaluate the creep behavior of G91 steel after different TMT and heat treatments. The minimum disk deflection rate was lower, and the time to rupture was longer for G91 after the TMT than with the conventional G91 heat treatment (AR). The improvement in creep rupture strength is attributed to the fine and homogeneous distribution of MX carbonitrides. The number density and average precipitate size of MX carbonitrides after the TMT are similar to the oxide particles in ODS steels. These latter steels possess high creep strength due to the high number density of oxides distributed in the matrix. Considering the MX carbonitrides as a substitute for oxides, 9Cr FM steels after the TMT are a potential replacement of ODS steels, which are fabricated by expensive powder metallurgy and mechanical alloying processing routes.
Creep failure: based on the results presented above and taking into account the different stages of the TMT, the loss of creep ductility that enhances the change in fracture mechanism would be promoted by the coarsening of M23C6 carbides at the vicinity of the prior austenite grains. The coarse M23C6 carbides located on prior austenite grain boundaries favor the nucleation of the cavities at the vicinity of the prior austenite grains. Besides, in the TMT samples, the high austenitization temperature produces an enormous prior austenite grain sizes with concomitant large grain boundary surfaces, facilitating an earlier formation of the critical crack length that causes the brittle fracture.
We acknowledge support for the publication fee by the CSIC Open Access Publication Support Initiative through its Unit of Information Resources for Research (URICI).
Motor imagery (MI) is the mental representation of a movement in the absence of any actual overt movement. It is a cognitive process creating specific motor actions within the working memory [1]. MI may be a beneficial tool to improve various motor functions for patients with stroke-induced motor deficits [2, 3, 4, 5]. Decline of motor evoked potential (MEP) amplitude, an index of corticospinal excitability obtained when transcranial magnetic stimulation is applied to the primary motor cortex, can be observed post-stroke [6]. Additionally, a significant reduction of spinal motor neuron excitability has been shown in the post-stroke acute phase [7]. Thus, corticospinal excitability, including that of spinal motor neurons, would be reduced post-stroke. Corticospinal excitability is considered to be an index of functional motor recovery [8], and immediate enrollment in rehabilitation for stimulation of corticospinal and spinal motor neuron excitability may be important to achieve better outcomes.
\nNeuroimaging studies show that MI activates motor-related brain regions, including the primary motor cortex, supplementary motor area, premotor area, prefrontal cortex, somatosensory area, parietal lobe, cingulate gyrus, cerebellum, and basal ganglia [9, 10]. Similarly, these regions have been shown as activated during motor execution [9, 10]. Furthermore, the MEP amplitude was shown as significantly increased during MI [11, 12, 13]. Thus, MI may stimulate the central nervous system. Various patterns of spinal motor neuron excitability have been observed during MI [14, 15, 16]. The F-wave is one of the indices of spinal motor neuron excitability. It is a compound action potential resulting from re-excitation (backfiring) at spinal anterior horn cells by an antidromic impulse following the distal electrical stimulation of motor nerve fibers [17, 18, 19]. The F-wave amplitude increases when the corticospinal descending volley collides with the antidromic peripheral volley [20]. Additionally, the F-wave is a reliable index of spinal motor neuron excitability, even when motor output is extremely low, as is the case during MI [21].
\nAs described previously in this chapter, stimulating spinal motor neuron excitability would improve motor function. Our final goal is to find the most beneficial approach by which MI can increase the spinal motor neuron excitability. In the following sections, we introduce our research on spinal motor neuron excitability under various MI conditions. At the end of the chapter, we discuss the application of MI to physical therapy from a neurophysiological perspective.
\nWe previously reported that the spinal motor neuron excitability increased significantly when participants performed MI of isometric thenar muscle activity under 50% maximal voluntary contraction (MVC) [22]. However, it was unclear whether the magnitude of the imagined muscle contraction strength affects the spinal motor neuron excitability. Therefore, we used F-wave measurements to investigate the spinal motor neuron excitability during MI of isometric thenar muscle activity under various imagined muscle contraction strengths. Specifically, we adopted the 10, 30, 50, 70, and 100% MVC for imagined muscle contraction strength [23, 24, 25].
\nWe conducted two experiments to assess the spinal motor neuron excitability during MI under different imagined muscle contraction strengths. Firstly, we measured the F-wave during MI under 10, 30, 50, and 70% MVC for 10 healthy volunteers (5 males, 5 females; mean age = 28.7 ± 4.5 years). Secondly, we measured the F-wave during MI under 50 and 100% MVC for 15 healthy volunteers (13 males, 2 females; mean age = 25.3 ± 5.0 years). All participants provided informed consent before study commencement. This research was approved by the Research Ethics Committee at Kansai University of Health Sciences. All recordings were conducted in accordance with the Declaration of Helsinki.
\nA Viking Quest electromyography (EMG) machine ver. 9.0 (Natus Medical Inc., USA) was used for the F-wave recording. A pair of silver disc electrodes (10 mm diameter, Natus Medical Inc., USA) were placed over left thenar eminence and base of the first dorsal metacarpal bone. The skin was cleaned with an abrasive gel, and then impedance was maintained below 5 kΩ. The F-wave was evoked from the left thenar muscles by delivering supramaximal electrical stimuli, 0.2 ms in duration and 0.5 Hz in frequency, to the median nerve at the left wrist. Supramaximal stimuli were determined 20% higher than the maximal stimulus intensity required to elicit the largest compound muscle action potential (M-wave). The sensitivity for the F-wave was set at 200 μV/division and a sweep of 5 ms/division. The bandwidth filter range was 20 Hz–3 kHz.
\nParticipants were placed in the supine posture on a bed and instructed to fix their eyes on the display of a pinch meter (Digital Indicator F304A, Unipulse Corp., Japan) throughout the F-wave recording. To determine the baseline of spinal motor neuron excitability, the F-wave was measured during relaxation for 1 min (rest). Thereafter, participants exerted isometric left thenar muscle contraction at 50% MVC (i.e., participants pressed the sensor of the pinch meter using their thumb and index finger at 50% MVC) for 1 min with visual feedback. For the MI trial, participants performed MI of isometric thenar muscle activity under 50% MVC for 1 min (50% MI). Immediately after the 50% MI trial, the F-wave was recorded during relaxation for 1 min (post). There were 30 supramaximal electrical stimuli delivered during each trial for the F-wave recording. The above process was defined as the MI at 50% MVC condition (50% MI condition). This protocol was repeated for conditions of 10, 30, 70, and 100% MI. Each condition was performed randomly on different days.
\nAll measured F-wave data were analyzed with respect to two parameters: persistence and the F/M amplitude ratio. The minimum F-wave peak-to-peak amplitude was 20 μV [26]. The persistence was represented as the percentage of detected F-wave responses out of 30 supramaximal electrical stimuli. It reflects the number of backfiring spinal anterior horn cells [17, 19]. The F/M amplitude ratio was obtained as the mean of the ratios of each detected F-wave response amplitude divided by the corresponding M-wave amplitude; it reflects the number of backfiring spinal anterior horn cells and individual spinal anterior horn cell excitability [19]. Therefore, the persistence and F/M amplitude ratio are considered indices of spinal motor neuron excitability.
\nBecause the Shapiro-Wilk test did not confirm the normality of the F-wave data, a nonparametric method was used for statistical analysis. The persistence and F/M amplitude ratio among the three trials (rest, MI, and post) under each MI condition (10, 30, 50, and 70% MI) were compared using the Friedman test and Scheffe’s post hoc test.
\nWe also calculated the relative value obtained by dividing the F-wave data during MI under the four MI conditions by that at rest and compared the results using the Friedman test. SPSS Statistics ver. 19 software (IBM Corp., USA) was used for statistical analysis. The threshold for statistical significance was set at p < 0.05.
\nThe persistence and F/M amplitude ratio among three trials (rest, MI, and post) under each MVC MI conditions were compared using the Friedman test and Scheffe’s post hoc test. The relative values between the two MI conditions were compared using the Wilcoxon signed rank test.
\nThe persistence during MI under all MI conditions was significantly higher than that at rest (10% MI vs. rest and 70% MI vs. rest, p < 0.01; 30% MI vs. rest and 50% MI vs. rest, p < 0.05) (\nTables 1\n–\n4\n). The F/M amplitude ratio during MI under 10, 30, and 50% MI conditions was significantly higher than that at rest (10% MI vs. rest and 50% MI vs. rest, p < 0.01; 30% MI vs. rest, p < 0.05) (\nTables 1\n\n–\n\n3\n). The F/M amplitude ratio during MI under the 70% MI condition tended to be more increased than that at rest (p = 0.082) (\nTable 4\n). The F/M amplitude ratio immediately after MI under all MI conditions was reduced to the rest level (\nTables 1\n\n–\n\n4\n).
\nThe F-wave under 10% MI condition.
The F-wave under 30% MI condition.
The F-wave under 50% MI condition.
The F-wave under 70% MI condition.
The relative values of the persistence and F/M amplitude ratio did not show significant differences among all MI conditions (\nTable 5\n).
\nRelative values of the F-wave under 10% MI, 30% MI, 50% MI, and 70% MI condition.
The persistence during MI under the 50% MI and 100% MI conditions was significantly higher than that at rest (50% MI vs. rest and 100% MI vs. rest, p < 0.01; \nTables 6\n and \n7\n), and the F/M amplitude ratio during MI under 50% and 100% MI conditions was significantly higher than that at rest (50% MI vs. rest and 100% MI vs. rest, p < 0.01; \nTables 6\n and \n7\n). The F/M amplitude ratio immediately after MI (at post) under the 50% and 100% MI conditions did not show any significant differences compared with that at rest (\nTables 6\n and \n7\n).
\nThe F-wave under 50% MI condition.
The F-wave under 100% MI condition.
The relative values of the persistence and F/M amplitude ratio did not show significant differences between two MI conditions (\nTable 8\n).
\nRelative values of the F-wave under 50% MI and 100% MI condition.
Both the persistence and the F/M amplitude ratio were significantly increased during MI under 10, 30, 50, 70, and 100% MVC. Previous research has demonstrated that the activation of various brain regions contributes to motor preparation and planning during MI [9, 10]. Thus, it is considered that the activation of the central nervous system that contributes to motor preparation and planning during MI is responsible for the observed increase in spinal motor neuron excitability via the descending pathways, such as the corticospinal and extrapyramidal tracts.
\nFurthermore, all participants in our previous studies performed MI while holding the sensor of the pinch meter. Mizuguchi et al. [27] reported that while holding an object, the corticospinal excitability during MI was modulated by a combination of tactile and proprioceptive inputs. Thus, it is plausible that holding the pinch meter sensor during MI caused tactile and proprioceptive perceptions to cooperatively increase the spinal motor neuron excitability along with the MI-activated pathways.
\nRelative values of the persistence and F/M amplitude were similar among all MI conditions. This result indicated that the magnitude of imagined muscle contraction strength may not affect spinal motor neuron excitability. Bonnet et al. [28] reported that the H-reflex amplitude during MI was similar between 2 and 10% MI conditions. Hale et al. [29] also reported that the H-reflex amplitude during MI of ankle plantar flexion was similar among five (i.e., 20, 40, 60, 80, and 100% MVC) MI conditions. Similarly, Aoyama and Kaneko [30] reported that the H-reflex amplitude during MI was similar between 50 and 100% MI conditions. MI is the mental representation of a movement in the absence of any overt movement [1]. The neural mechanism that inhibits actual movement and muscle contraction during MI may be involved in this result. Park et al. [31] reported that the MEP amplitude during MI was similar among all six (i.e., 10, 20, 30, 40, 50, and 60% MVC) MI conditions. Furthermore, the magnitude of primary motor cortex activity during MI did not correlate with that of the imagined muscle contraction strength, although the activities of the supplementary motor and premotor area during MI were strongly correlated with it [32]. The supplementary motor and premotor areas play crucial roles in larger force generation [33], motor planning, preparation, and inhibition [34, 35]. Thus, these areas may inhibit the actual muscle contraction depending on the magnitude of the muscle contraction strength. These areas are also directly connected to the primary motor cortex, and inhibitory inputs from them may suppress additional primary motor cortex excitation conferred by MI with high imagined contraction strengths. Therefore, the degree of spinal motor neuron excitability during MI at various imagined muscle contraction strengths may be modulated by both excitatory and inhibitory inputs from the central nervous system.
\nOur previous research has shown significant facilitation of the spinal motor neuron excitability during MI of isometric thenar muscle activity. The imagined muscle contraction strength may not be affected by the spinal motor neuron excitability. Thus, MI of isometric thenar muscle activity under slight MVC (i.e., 10% MVC) could substantially facilitate the spinal motor neuron excitability.
\nWe previously reported that MI can increase the spinal motor neuron excitability and that the magnitude of imagined muscle contraction strength may not affect spinal motor neuron excitability [23, 24, 25]. In these studies, the duration of each MI session was 1 min. Driskell et al. [36] suggested that longer MI sessions do not always prove beneficial; they recommended a duration of approximately 20 min for an MI training session. Another study suggested that MI for 10–15 min elicited the most significant effect on performance [37]. Moreover, Twinning et al. [38] suggested that 5 min is the temporal limit beyond which it is difficult to concentrate and perform MI. As described previously, stimulation of spinal motor neuron excitability may be important for post-stroke rehabilitation; however, time-dependent changes in spinal motor neuron excitability during MI have not yet been investigated. Additionally, MI ability has a significant effect on brain activation [39] and corticospinal excitability [40]. In this study, we used F-waves to investigate whether the duration of MI and MI ability affects the spinal motor neuron excitability [41].
\nEleven healthy volunteers participated in this research (8 males, 3 females, mean age = 26.4 ± 6.0 years). All participants gave written informed consent before study commencement. The study was approved by the Research Ethics Committee at the Graduate School of Kansai University of Health Sciences. All recordings were conducted in accordance with the Declaration of Helsinki.
\nThe environment and F-wave recording conditions were set as previously described [23, 24, 25]. For the rest trial (rest), the F-wave was measured during relaxation for 1 min. Subsequently, participants learned the isometric thenar muscle activity at 50% MVC for 1 min with visual feedback. For the MI trial, participants performed the MI of isometric thenar muscle activity under 50% MVC for 5 min. F-waves were measured at 1, 3, and 5 min after the beginning of MI (1-, 3-, and 5-min MI, respectively). Immediately after MI, the F-wave was measured again during relaxation (post). After F-wave recordings, participants were asked to evaluate their vividness of MI, how vividly they could imagine isometric thenar muscle activity at 50% MVC, at 1, 3, and 5-min using a seven-point Likert scale ranging from 1 (very difficult to perform MI vividly) to 7 (very easy to perform MI vividly).
\nBackground electromyography (EMG) was recorded using telemetry EMG (MQ-8, Kissei Comtec Co., Ltd., Japan) and EMG recording software (Vital Recorder 2, Kissei Comtec Co., Ltd.). Surface EMG signals were recorded for 5 min from the left thenar muscles to confirm no muscle contractions during MI. A pair of disposable Ag/AgCl electrodes (Blue Sensor N-00-S, Ambu A/S, Denmark) were placed over the muscle surface with an inter-electrode distance of 20 mm. EMG signals were recorded at rest; at 1, 3, and 5 min of MI; and post-trial. The recorded EMG data were analyzed using a multi-purpose biological information analysis system (BIMUTAS-Video, Kissei Comtec Co., Ltd.) after analog to digital conversion at a sampling frequency of 1 kHz. The root mean square values of the EMG data in each trial were then calculated.
\nWe used a nonparametric method because the normality of F-wave data was not confirmed using the Shapiro-Wilk test. The persistence and F/M amplitude ratio among five trials (rest, 1, 3, 5-min MI, and post, respectively) were compared using the Friedman test and Scheffe’s post hoc test. The rating scores of MI vividness at 1, 3, and 5-min MI were compared using the Friedman test and Scheffe’s post hoc test. The SPSS statistics ver. 19 (IBM Corp., USA) was used for statistical analysis. The threshold for statistical significance was set at p = 0.05.
\nIn \nFigure 1\n it is possible to verify that persistence and F/M amplitude ratio were significantly facilitated until 3 min from the beginning of MI task.
\nChanges in persistence and F/M amplitude ratio during MI for 5 min (*p < 0.05, **p < 0.01). The persistence at 1 and 3 min MI was significantly higher than that at rest. The F/M amplitude ratio at 1 and 3 min MI was significantly higher than that at rest. Additionally, the F/M amplitude ratio at 5 min MI was significantly smaller than that at 1 and 3 min MI.
The score of MI vividness at 5-min MI was significantly decreased compared to 1-min MI (*p < 0.05; \nTable 9\n).
\nRating scores of MI vividness at 1-min MI, 3-min MI, and 5-min MI.
There were no significant differences in the RMS data among five trials, and thus there was no measurable muscle activity during MI for 5 min.
\nBoth persistence and F/M amplitude ratio were significantly higher until 3 min from the beginning of MI. This result may suggest that participants could perform MI for at least 3 min without much difficulty. However, there were no significant differences in persistence and F/M amplitude ratio between the at rest and 5-min MI conditions. Additionally, the F/M amplitude ratio for the 5-min MI conditions was significantly lower than that for 1- and 3-min conditions. These findings may be due to mental fatigue; in one study, mental fatigue was found to have altered the maximal force production of the elbow flexor [42]. It also made it difficult for participants to maintain their focus on imagined movement [43]. Furthermore, repetitive MI of a handgrip movement decreased the MEP amplitude more than that at rest [44]. Thus, mental fatigue caused by sustained mental activity may have induced a decline of the spinal motor neuron excitability.
\nFurthermore, a decline of spinal motor neuron excitability can be also explained by MI habituation. MI is closely related to attentional processing [45]. Brain activation decreases by habituation after performing a cognitive motor task for 10 min. Furthermore, the corticospinal excitability was also decreased by habituation [46]. Specifically, brain activity showed an increase at 2 min before the onset of the task; however, after 4–6 min, activity decreased. Additionally, at the spinal level, the T-reflex amplitude, another index of the spinal motor neuron excitability, was significantly decreased due to habituation following sustained mental work for 20 min [47]. Our results also seemed to indicate that habituation to MI might occur approximately 4 min after its initiation and suggested that longer excitation times during MI might not be required for habituation of the central nervous system and spinal motor neurons.
\nFinally, practice time and MI ability were considered as possible factors affecting spinal motor neuron excitability. Regarding clinical use of MI for motor skill learning, Twining et al. [38] indicated that participants found it difficult to concentrate and perform MI for more than 5 min. Mental chronometry measured similar times for actual performance and MI [48]. Specifically, participants experienced difficulties in performing MI accurately beyond the practice time. In our study, the practice time for the motor task was only 1 min; thus, 1 min of practice time may be insufficient to continue performing MI for 5 min. Indeed, the vividness of MI tended to decrease with MI time. Furthermore, the vividness of MI at 5 min post MI initiation was significantly decreased relative to that at 1 min post MI initiation.
\nThe persistence and F/M amplitudes at 1- and 3-min MI were significantly increased; however, the persistence and the F/M amplitude ratio at 5-min MI were reduced to rest levels. Thus, MI for 1–3 min may positively affect the spinal motor neuron excitability. In physical therapy, the duration of MI should be considered. As described in the Discussion section, matching the time of task practice to that of MI might be important. However, in this study, we did not investigate time-dependent changes of the spinal motor neuron excitability after motor learning for 5 min. Therefore, further research is required to resolve this issue. A limitation of this research is that differences in the brain activity during MI under 10, 30, 50, 70, and 100% MVC were not evaluated. Further study would be required to resolve this issue.
\nPrevious research has demonstrated that MI increases the spinal motor neuron excitability and that the magnitude of imagined contraction strength may not affect it [23, 24, 25]. Additionally, the duration of MI should be considered in physical therapy [41].
\nMI includes various components of perception that can be associated with actual movement [49], which is why the effects of MI may differ depending on the choice of sensory modality. Here, we used F-wave and MI ability to investigate whether the choice of imagery strategy affects the spinal motor neuron excitability [50].
\nFourteen healthy volunteers participated in this research (10 males, 4 females, mean age = 23.4 ± 4.8 years). All participants gave written informed consent before study commencement. The study was approved by the Research Ethics Committee at the Graduate School of Kansai University of Health Sciences. All recordings were conducted in accordance with the Declaration of Helsinki.
\nThe environment and F-wave recording conditions were set as previously described [23, 24, 25, 41]. To determine the baseline of the spinal motor neuron excitability, the F-wave was measured during relaxation for 1 min (rest). Subsequently, participants exerted isometric left thenar muscle contraction at 50% MVC for 1 min with visual feedback. Simultaneously, participants were instructed to learn the two imagery strategies: somatosensory (tactile and pressure perception of thumb finger pulp during pressing of the sensor of the pinch meter) and kinesthetic (thenar muscle contraction during pressing of the sensor of the pinch meter at 50% MVC). After learning each imagery strategy, participants performed somatosensory imagery (SI), kinesthetic imagery (KI), and combined somatosensory and kinesthetic imagery (SKI) randomly for 1 min. In SKI trial, participants performed kinesthetic and somatosensory imagery simultaneously. After all the F-wave recording, participants were asked to evaluate difficulty of each imagery strategies by using a 5-point Likert scale, ranging from 1 (very hard to image vividly) to 5 (very easy to image vividly).
\nBackground electromyography (EMG) was recorded during rest and three imagery trials.
\nWe used a nonparametric method because the normality of F-wave data was not confirmed using the Shapiro-Wilk test. The persistence and F/M amplitude ratio among four trials (rest, SI, KI, and SKI, respectively) were compared using the Friedman test and Scheffe’s post hoc test. The rating scores of each imagery strategies (SI, KI, and SKI, respectively) were compared using the Friedman test and Scheffe’s post hoc test. The background EMG data were compared using the Friedman test. The SPSS statistics ver. 19 (IBM Corp., USA) was used for statistical analysis. The threshold for statistical significance was set at p = 0.05.
\nThe persistence during SI and KI were significantly higher than that at rest (**p < 0.01; \nFigure 2\n). The persistence during SKI was tended to be increased than that at rest (p = 0.097; \nFigure 2\n). The F/M amplitude ratio during KI was significantly higher than that at rest (*p < 0.05; \nFigure 2\n).
\nChanges in persistence and F/M amplitude ratio during SI, KI, and SKI (*p < 0.05, **p < 0.01). The persistence during SI and KI was significantly higher than that at rest. The persistence during SKI was tended to be increased than that at rest. The F/M amplitude ratio during KI was significantly higher than that at rest.
The rating score of SKI vividness was significantly lower than that at rest (*p < 0.05; \nTable 10\n).
\nRating scores of MI vividness during SI, KI, and SKI.
There were no significant differences in the background EMG data among four trials (rest, SI, KI and SKI, respectively), and thus there was no measurable muscle contraction during three imagery trials.
\nBoth persistence and F/M amplitude ratio were significantly higher than the corresponding at-rest values. Previous neurophysiological studies reported that various regions of the brain related to motor functions were activated [9, 10] and that the MEP amplitude was significantly increased during KI [11, 13]. Thus, it seems that the central nervous system can better stimulate spinal motor neuron excitability via the descending pathways.
\nThe persistence during SI was significantly higher than that at rest. This result was unexpected. We previously hypothesized that the spinal motor neuron excitability would remain unchanged because there are no motor-related factors in SI. Furthermore, there are no previous reports of tactile and proprioceptive perception SI increasing the corticospinal excitability including that of the primary motor cortex. Thus, it may be difficult to increase the spinal motor neuron excitability by SI. However, it is possible for SI to include kinesthetic components. Participants in this research were asked to imagine tactile and pressure perception while holding the pinch meter sensor between their thumb and index finger. Thus, they might have unintentionally imagined tactile and pressure perception along with thenar muscle activity.
\nThe persistence during SKI tended to be increased as compared to that at rest. The rating score of SKI vividness was the lowest among the three imagery strategies. These results indicate that participants may not be able to perform SKI as vividly as the other two strategies. In this study, participants were required to pay attention to kinesthetic and somatosensory perceptions simultaneously. The decline in the amount of attention that can be allocated to each imagery strategy may have made it difficult for the participants to perform SKI vividly. Indeed, there was a positive correlation between the corticospinal excitability and MI vividness [40].
\nFrom the result of this research, KI may be a more effective imagery strategy, which can increase the spinal motor neuron excitability. Thus, the imagery strategy should be considered in physical therapy. Also, the spinal motor neuron excitability during SI was significantly increased. However, the mechanism that SI increases spinal motor neuron excitability is unclear. As a limitation of this research, we did not investigate brain activity during SI. Further research will be required to resolve this limitation.
\nOur research indicates that MI of the isometric thenar muscle activity can increase spinal motor neuron excitability [22, 23, 24, 25, 41, 50]. After a stroke or a spinal cord injury, the excitability of the central nervous system decreases due to various factors, including the damage of neural substrates, loss of sensory inputs, and disuse of affected limbs [51]. Additionally, corticospinal excitability decreases following a decline in the size and number of corticospinal neurons [52]. Furthermore, decline of spinal motor neuron excitability was shown in the post-stroke acute phase [7, 53]. Thus, it may be important to stimulate corticospinal excitability, including that of the spinal motor neurons, as soon as possible. Patients in an early postoperative and post-stroke stage have difficulties in performing physical activities. However, considering the characteristics of MI, it can be a beneficial method to stimulate spinal motor neuron excitability without any overt movement and muscle contraction.
\nFurthermore, MI can improve not only the spinal motor neuron excitability but also various motor functions. Yue et al. [54] indicated that MI under 100% MVC for 4 weeks can increase the muscle strength of little finger abduction. Additionally, Sidaway et al. [55] indicated that MI under 100% MVC for 4 weeks can increase muscle strength of ankle dorsiflexion. About these results, Grosprêtre et al. [56] considered that MI may strengthen brain-to-muscle communication, including the enhanced recruitment of spinal motor neurons and involvement of the descending command. Although other groups [54, 55] adopted maximal imagined muscle contraction strengths for MI training, our results [23, 24, 25] revealed that the magnitude of imagined muscle contraction strength did not affect spinal motor neuron excitability. Thus, low (i.e., 10% MVC) imagined muscle contraction strengths might be sufficient for stimulation of spinal motor neuron excitability and muscle strength. Our research [25, 41, 50] also revealed that kinesthetic imagery can better stimulate spinal motor neuron excitability and that spinal motor neuron excitability remained higher than the at-rest value until 3 min after MI initiation. Therefore, to increase the effects of MI, kinesthetic perception should be chosen as the imagery strategy. Additionally, the duration of each MI session should be less than 3 min.
\nIn conclusion, MI can increase the spinal motor neuron excitability, and its effect would be changed depending on the duration and strategy of imagery. Thus, the duration and strategy of imagery should be considered in clinical settings.
\nThe author would like to thank Prof. Toshiaki Suzuki from Graduate school of Kansai University of Health Sciences for helpful comments on this manuscript.
\nThere is no conflict of interest.
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\n\nMetadata for all publications is also automatically deposited in IntechOpen's OAI repository, making them available through the Open Access Infrastructure for Research in Europe's (OpenAIRE) search interface further establishing our compliance.
\n\nIn other words, publishing with IntechOpen guarantees compliance.
\n\nRead more about Open Access in Horizon 2020 here.
\n\nWhich scientific publication to choose?
\n\nWhen choosing a publication, Horizon 2020 grant recipients are encouraged to provide open access to various types of scientific publications including monographs, edited books and conference proceedings.
\n\nIntechOpen publishes all of the aforementioned formats in compliance with the requirements and criteria established by the European Commission for the Horizon 2020 Program.
\n\nAuthors requiring additional information are welcome to send their inquiries to funders@intechopen.com
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