Vol. 26 No. 1 2024 3 EDITORIAL COUNCIL EDITORIAL BOARD EDITOR-IN-CHIEF: Anatoliy A. Bataev, D.Sc. (Engineering), Professor, Rector, Novosibirsk State Technical University, Novosibirsk, Russian Federation DEPUTIES EDITOR-IN-CHIEF: Vladimir V. Ivancivsky, D.Sc. (Engineering), Associate Professor, Department of Industrial Machinery Design, Novosibirsk State Technical University, Novosibirsk, Russian Federation Vadim Y. Skeeba, Ph.D. (Engineering), Associate Professor, Department of Industrial Machinery Design, Novosibirsk State Technical University, Novosibirsk, Russian Federation Editor of the English translation: Elena A. Lozhkina, Ph.D. (Engineering), Department of Material Science in Mechanical Engineering, Novosibirsk State Technical University, Novosibirsk, Russian Federation The journal is issued since 1999 Publication frequency – 4 numbers a year Data on the journal are published in «Ulrich's Periodical Directory» Journal “Obrabotka Metallov” (“Metal Working and Material Science”) has been Indexed in Clarivate Analytics Services. Novosibirsk State Technical University, Prospekt K. Marksa, 20, Novosibirsk, 630073, Russia Tel.: +7 (383) 346-17-75 http://journals.nstu.ru/obrabotka_metallov E-mail: metal_working@mail.ru; metal_working@corp.nstu.ru Journal “Obrabotka Metallov – Metal Working and Material Science” is indexed in the world's largest abstracting bibliographic and scientometric databases Web of Science and Scopus. Journal “Obrabotka Metallov” (“Metal Working & Material Science”) has entered into an electronic licensing relationship with EBSCO Publishing, the world's leading aggregator of full text journals, magazines and eBooks. The full text of JOURNAL can be found in the EBSCOhost™ databases.
OBRABOTKAMETALLOV Vol. 26 No. 1 2024 4 EDITORIAL COUNCIL EDITORIAL COUNCIL CHAIRMAN: Nikolai V. Pustovoy, D.Sc. (Engineering), Professor, President, Novosibirsk State Technical University, Novosibirsk, Russian Federation MEMBERS: The Federative Republic of Brazil: Alberto Moreira Jorge Junior, Dr.-Ing., Full Professor; Federal University of São Carlos, São Carlos The Federal Republic of Germany: Moniko Greif, Dr.-Ing., Professor, Hochschule RheinMain University of Applied Sciences, Russelsheim Florian Nürnberger, Dr.-Ing., Chief Engineer and Head of the Department “Technology of Materials”, Leibniz Universität Hannover, Garbsen; Thomas Hassel, Dr.-Ing., Head of Underwater Technology Center Hanover, Leibniz Universität Hannover, Garbsen The Spain: Andrey L. Chuvilin, Ph.D. (Physics and Mathematics), Ikerbasque Research Professor, Head of Electron Microscopy Laboratory “CIC nanoGUNE”, San Sebastian The Republic of Belarus: Fyodor I. Panteleenko, D.Sc. (Engineering), Professor, First Vice-Rector, Corresponding Member of National Academy of Sciences of Belarus, Belarusian National Technical University, Minsk The Ukraine: Sergiy V. Kovalevskyy, D.Sc. (Engineering), Professor, Vice Rector for Research and Academic Aff airs, Donbass State Engineering Academy, Kramatorsk The Russian Federation: Vladimir G. Atapin, D.Sc. (Engineering), Professor, Novosibirsk State Technical University, Novosibirsk; Victor P. Balkov, Deputy general director, Research and Development Tooling Institute “VNIIINSTRUMENT”, Moscow; Vladimir A. Bataev, D.Sc. (Engineering), Professor, Novosibirsk State Technical University, Novosibirsk; Vladimir G. Burov, D.Sc. (Engineering), Professor, Novosibirsk State Technical University, Novosibirsk; Aleksandr N. Korotkov, D.Sc. (Engineering), Professor, Kuzbass State Technical University, Kemerovo; Dmitry V. Lobanov, D.Sc. (Engineering), Associate Professor, I.N. Ulianov Chuvash State University, Cheboksary; Aleksey V. Makarov, D.Sc. (Engineering), Corresponding Member of RAS, Head of division, Head of laboratory (Laboratory of Mechanical Properties) M.N. Miheev Institute of Metal Physics, Russian Academy of Sciences (Ural Branch), Yekaterinburg; Aleksandr G. Ovcharenko, D.Sc. (Engineering), Professor, Biysk Technological Institute, Biysk; Yuriy N. Saraev, D.Sc. (Engineering), Professor, V.P. Larionov Institute of the Physical-Technical Problems of the North of the Siberian Branch of the RAS, Yakutsk; Alexander S. Yanyushkin, D.Sc. (Engineering), Professor, I.N. Ulianov Chuvash State University, Cheboksary
Vol. 26 No. 1 2024 5 CONTENTS OBRABOTKAMETALLOV TECHNOLOGY Kuts V.V., Oleshitsky A.V., Grechukhin A.N., Grigorov I.Y. Investigation of changes in geometrical parameters of GMAW surfaced specimens under the infl uence of longitudinal magnetic fi eld on electric arc....................................... 6 Saprykina N.А., Chebodaeva V.V., Saprykin A.А., Sharkeev Y.P., Ibragimov E.А., Guseva T.S. Optimization of selective laser melting modes of powder composition of the AlSiMg system................................................................. 22 Gubin D.S., Kisel’ A.G. Features of calculating the cutting temperature during high-speed milling of aluminum alloys without the use of cutting fl uid............................................................................................................................................. 38 EQUIPMENT. INSTRUMENTS Borisov M.A., Lobanov D.V., Zvorygin A.S., Skeeba V.Y. Adaptation of the CNC system of the machine to the conditions of combined processing...................................................................................................................................... 55 Nosenko V.A., Bagaiskov Y.S., Mirocedi A.E., GorbunovA.S. Elastic hones for polishing tooth profi les of heat-treated spur wheels for special applications..................................................................................................................................... 66 Podgornyj Y.I., Skeeba V.Y., Martynova T.G., Lobanov D.V., Martyushev N.V., Papko S.S., Rozhnov E.E., Yulusov I.S. Synthesis of the heddle drive mechanism....................................................................................................... 80 MATERIAL SCIENCE Ragazin A.A., Aryshenskii V.Y., Konovalov S.V., Aryshenskii E.V., Bakhtegareev I.D. Study of the eff ect of hafnium and erbium content on the formation of microstructure in aluminium alloy 1590 cast into a copper chill mold............................................................................................................................................................................ 99 Zorin I.A., Aryshenskii E.V., Drits A.M., Konovalov S.V. Study of evolution of microstructure and mechanical properties in aluminum alloy 1570 with the addition of 0.5 % hafnium........................................................................... 113 Karlina Y.I., Kononenko R.V., Ivantsivsky V.V., Popov M.A., Deryugin F.F., Byankin V.E. Relationship between microstructure and impact toughness of weld metals in pipe high-strength low-alloy steels (research review)..................... 129 Patil N.G., Saraf A.R., Kulkarni A.P Semi empirical modeling of cutting temperature and surface roughness in turning of engineering materials with TiAlN coated carbide tool................................................................................. 155 Sawant D., Bulakh R., Jatti V., Chinchanikar S., Mishra A., Sefene E.M. Investigation on the electrical discharge machining of cryogenic treated beryllium copper (BeCu) alloys........................................................................................ 175 Karlina A.I., Kondratiev V.V., Sysoev I.A., Kolosov A.D., Konstantinova M.V., Guseva E.A. Study of the eff ect of a combined modifi er from silicon production waste on the properties of gray cast iron................................................. 194 EDITORIALMATERIALS 212 FOUNDERS MATERIALS 223 CONTENTS
OBRABOTKAMETALLOV MATERIAL SCIENCE Vol. 26 No. 1 2024 Relationship between microstructure and impact toughness of weld metals in pipe high-strength low-alloy steels (research review) Yulia Karlina 1, a, *, Roman Kononenko 2, b, Vladimir Ivancivsky 3, c, Maksim Popov 2, d, Fedor Derjugin 2, e, Vladislav Byankin 2, f 1 National Research Moscow State University of Civil Engineering, 26 Yaroslavskoe Shosse, Moscow, 129337, Russian Federation 2 Irkutsk National Research Technical University, 83 Lermontova str., Irkutsk, 664074, Russian Federation 3 Novosibirsk State Technical University, 20 Prospekt K. Marksa, Novosibirsk, 630073, Russian Federation a https://orcid.org/0000-0001-6519-561X, jul.karlina@gmail.com; b https://orcid.org/0009-0001-5900-065X, istu_politeh@mail.ru; c https://orcid.org/0000-0001-9244-225X, ivancivskij@corp.nstu.ru; d https://orcid.org/0000-0003-2387-9620, popovma.kvantum@gmail.com; e https://orcid.org/0009-0004-4677-3970, deryugin040301@yandex.ru; f https://orcid.org/0009-0007-0488-2724, borck3420@gmail.com Obrabotka metallov - Metal Working and Material Science Journal homepage: http://journals.nstu.ru/obrabotka_metallov Obrabotka metallov (tekhnologiya, oborudovanie, instrumenty) = Metal Working and Material Science. 2024 vol. 26 no. 1 pp. 129–154 ISSN: 1994-6309 (print) / 2541-819X (online) DOI: 10.17212/1994-6309-2024-26.1-129-154 ART I CLE I NFO Article history: Received: 19 September 2023 Revised: 21 October 2023 Accepted: 16 January 2024 Available online: 15 March 2024 Keywords: Steel Ferrite Perlite Beinite Martensite Impact toughness Fracture Hybrid laser welding Standards Acknowledgements Research was partially conducted at core facility “Structure, mechanical and physical properties of materials”. ABSTRACT Introduction. The modern pipeline industry requires the development of materials of high strength and toughness for the production of steels for oil and gas pipelines. Changes in steel production and rolling technologies have become a challenge for developers of welding materials and joining technologies. This problem is more critical for strength levels above 830 MPa, where there are no special rules for the approval of welding consumables. Research methods. The failure of stainless steel pipeline welds is becoming a serious problem in the pipeline industry. Multiphase microstructures containing acicular ferrite or an acicular ferrite-dominated phase exhibit good complex properties in HSLA steels. This paper focuses on the results obtained using modern methods of scanning electron microscopy for microstructural analysis, backscattered electrons (BSE) for electron channel contrast imaging (ECCI) and orientation microscopy based on electron backscatter diff raction (ORM), as well as characteristic X-rays for compositional analysis using X-beam spectroscopy (XEDS) and secondary electrons (SE) to observe surface morphology. Results and discussion. This paper analyzes the characteristics of the microstructure of the weld and its relationship with impact toughness. It is shown that predicting impact toughness based on the microstructural characteristics of steel weld metals is complicated due to the large number of parameters involved. This requires an optimal microstructure of the steel. Satisfactory microstructure depends on several factors, such as chemical composition, hot work processing, and accelerated cooling. Alloying elements have a complex eff ect on the properties of steel, and alloying additives commonly added to the steel composition include Mn, Mo, Ti, Nb and V. From a metallurgical point of view, the choice of alloying elements and the metallurgical process can greatly infl uence the resulting microstructure. A longer cooling time tend to improve the toughness and reduce the mechanical strength of weld deposits on high-strength steels. Welding thermal cycles cause signifi cant changes in the mechanical properties of the base material. The analysis showed that impact toughness strongly depends on the microstructure of the multi-pass weld of the material under study, which contains several sources of heterogeneity, such as interdendritic segregation, and the eff ective grain size can also be a signifi cant factor explaining large deviations in local impact toughness values. Acicular ferrite nucleated in intragranular inclusions has been shown to produce a fi ne-grained interlocking arrangement of ferrite plates providing high tensile strength and excellent toughness, and is therefore a desirable microstructural constituent in C-Mn steel weld metals. At the same time, discussion regarding the relationship between acicular ferrite and toughness is very complex and still open at present. Relating impact toughness to acicular ferrite, taking into account the top bead, is not a reliable procedure, even for single-pass deposit welding. Impact strength depends on several factors, and the strong eff ect of acicular ferrite is generally recognized due to its fi ne-grained interlocking structure, which prevents the propagation of brittle cracks by cleavage. The large-angle boundaries and high dislocation density of acicular ferrite provide high strength and toughness. However, for the same amount of acicular ferrite, diff erent viscosity values may be observed depending on the content of microalloying elements in the steel. An analysis of the results of various studies showed that other factors also aff ect the impact strength. For example, microphases present along the Charpy-V notch are critical for the toughness of weld metals. The combination of OM, SEM and EBSD techniques provides an interesting method for metallographic investigation of the refi ned metal microstructure of stainless steel pipeline welds. Conclusion. This review reports the most representative study regarding the microstructural factor in the weld of pipe steels. It includes a summary of the most important process variables, material properties, regulatory guidelines, and microstructure characteristics and mechanical properties of the joints. This review is intended to benefi t readers from a variety of backgrounds, from non-welding or materials scientists to various industrial application specialists and researchers. For citation: Karlina Y.I., Kononenko R.V., Ivancivsky V.V., Popov M.A., Derjugin F.F., Byankin V.E. Relationship between microstructure and impact toughness of weld metals in pipe high-strength low-alloy steels (research review). Obrabotka metallov (tekhnologiya, oborudovanie, instrumenty) = Metal Working and Material Science, 2024, vol. 26, no. 1, pp. 129–154. DOI: 10.17212/1994-6309-2024-26.1-129-154. (In Russian). ______ * Corresponding author Karlina Yulia I., Ph.D. (Engineering), Research Associate National Research Moscow State Construction University, Yaroslavskoe shosse, 26, 129337, Moscow, Russian Federation Tel.: +7 914 879-85-05, e-mail: jul.karlina@gmail.com
OBRABOTKAMETALLOV MATERIAL SCIENCE Vol. 26 No. 1 2024 Introduction In the review [1], the features of the chemical composition of pipe steels, welding methods, and regulatory documents regulating mechanical properties are considered. In this paper we will consider the characteristics of the microstructure of welded joints. Increasing the yield strength is known to increase the loading capacity and reduce the transportation costs. Thus, high strength combined with high toughness and formability are the main requirements in the steel industry for pipelines [2–10]. The addition of micro-alloying elements such as Nb, V, Ti and Mo, coupled with advanced thermo-mechanical control process (TMCP) technology, can provide an excellent combination of strength and toughness [2, 3]. Microalloying elements such as Ti and Nb form fi nely dispersed carbide and carbonitride precipitates during TMCP of high-quality pipeline steels, which increase the strength of the steel. It has been established that fairly homogeneous dispersed particles containing Nb, Ti and V eff ectively inhibit the growth of austenite grains [11–15]. Besides, the additions of Mo, Nb and Cu contributed to the formation of a bainite microstructure [11–16]. The eff ect of carbide size on fracture may be indirectly related to grain size. The authors of [3, 11, 12] noted that the largest size of carbides in the microstructure is proportional to the size of the ferrite grain in annealed or normalized steels. Grain size is important even when cracks are initiated by pearlite particles or colonies, [11, 12] because the grains around the fracture source can control crack propagation [1–3]. Larger grains, if present around the source of the spall, encourage the nucleated crack to grow beyond the critical size required for unstable propagation before it can be blocked by the grain boundary. As a result, failure occurs at a lower stress than required when smaller grains are present around the start of the fracture. Observations such as the presence of non-propagating ferrite grain-sized cracks on the fracture surface [11], large cleavage facets at the crack nucleation (larger than the average facet size) [12–15], and better correlation between fracture stress and largest grain size (and not the average grain size) in fractured ferritic-pearlite steel specimens [17–25] is important in the initiation and propagation of cleavage cracks. At the same time, it should be understood that within the volume of a structural material, spatial inhomogeneities can arise in various forms, such as a non-uniform distribution of non-metallic inclusions and precipitates, a spatial distribution of pearlite and ferrite, a mixed (fi ne- and coarse-grained) granular structure (or crystallographic texture) [1–3]. The authors [3, 11, 12, 24, 25] concluded that spatial heterogeneity in any form can lead to a wider than usual scatter of fracture toughness results, depending on the local microstructure sampled at the “critical distance” (at where the local tensile stress exceeds the cleavage stress). Fracture stress [25] in front of the notch root. The grain size in steels can be irregular, and in some Nb-V steel plates subjected to thermo-mechanical control process (TMCP), a bimodal ferrite grain size distribution has been reported (coarse grains present in a matrix of fi ne grains) [11]. Therefore, depending on whether the grains are large or small at the root of the notch, the fracture stress values for a bimodal ferrite structure may diff er. Understanding the spread of Charpy energy values for steels after TMCP is very important from an industrial point of view. However, it is scientifi cally diffi cult to study the eff ect of particle size distribution on impact strength using Charpy tests. Charpy tests often produce complex fracture surfaces that make it diffi cult to identify the original location of the onset of cleavage [11, 25–28]. For example, works [11, 12] have shown that in a blunt notch test, if a coarse grain band is present in the active area just before the root of the notch, the coarse grains initiate spalling, which results in low shear failure stress. However, if large grains are absent at the root of the notch, small grains initiate spalling and fracture stress values are higher. Similarly, in the Charpy impact transition (IT) region, the magnitude of the plastic fracture area depends on the location of the coarse grain band relative to the root of the notch. If the coarse grain band is located close to the base of the cut, cleavage failure begins at that location, resulting in low impact energy. However, if the coarse grain strip is located far from the base of the notch, a ductile crack will propagate fi rst, absorbing higher impact energy. In addition to the above, the works [3, 11, 12, 15–19] show that the addition of a large amount of micro-alloying elements poses a serious problem for the weldability of pipeline steel due to the increased equivalent carbon content (Ceq according to the Russian standard), especially such elements as Ni, V, Cr, Mo and Cu [2, 4, 11–28].
OBRABOTKAMETALLOV MATERIAL SCIENCE Vol. 26 No. 1 2024 Research methods Predicting impact toughness based on the microstructural characteristics of weld metals is diffi cult due to the large number of parameters involved [1, 11–18]. The common practice of relating impact toughness to the microstructure of the last bead of a multi-pass weld turned out to be unsatisfactory since the amount of acicular ferrite, the most desirable component, may not always make the main contribution to the impact toughness [20–32]. Parameters such as the recrystallized fraction, the presence of micro-phases and inclusions can also play an important role [32–36, 37–48]. Thus, in order to take into account the infl uence of all these parameters, the method [38, 39] proposed by the International Institute of Welding (IIW) is not comprehensive enough and therefore additional methods are needed. This situation is more relevant for high-strength steel weldments where very fi ne microstructures cannot be clearly identifi ed, resulting in incorrect microstructure identifi cation. The use of scanning electron microscopy as an auxiliary method to optical microscopy has been successfully used for many decades to study C-Mn and low-alloy metal welds, mainly in the assessment of refi ned microstructure. Recently [49–61], in addition to the previously mentioned methods, electron backscatter diff raction (EBSD) has also been used to provide a more effi cient analytical procedure. This method, which provides valuable grain boundary information, is useful for refi ned microstructures to confi rm constituents such as acicular ferrite, bainite and martensite. The mechanical properties of high-strength low-alloy pipe steels largely depend on its complex microstructure. However, the precise quantitative infl uence of individual microstructural elements (e.g., dislocations, grain boundaries, phase boundaries, volume fractions of the corresponding microstructure components, phase types, dispersion and shape of martensite islands, etc.) [2, 3, 11] is usually not easy to measure with traditional optical microscopymethods. Thus, it is a general question how to obtain quantitative values of the types and quantities of these diff erent microstructural components and its topological features. Various electron diff raction techniques, mainly used in scanning electron microscopy (SEM), are capable of providing comprehensive answers to these questions. Modern scanning electron microscopes with thermal fi eld emission guns, various sensitive detectors and fl exible stages are extremely versatile tools for detailed and quantitative analysis of the microstructure of bulk materials’ specimens with high resolution, with large statistics, in 2D and 3D, as well as provide the opportunity to implement various types of fi eld observations. The most important signals to be detected for microstructural analysis are backscattered electrons (BSE) for electron channel contrast imaging (ECCI) and orientation microscopy based on electron backscatter diff raction (ORM), as well as characteristic X-rays for compositional analysis using X-beam spectroscopy (XEDS) and secondary electron (SE) to observe surface morphology. The purpose of the work is to evaluate the various microstructures of metal welds of C-Mn and highstrength steels based on the analysis of various studies carried out by optical microscopy, scanning electron microscopy and EBSD methods, taking into account the infl uence of recrystallization in multi-pass welds, microstructural components, micro-phases, and inclusion. The objective of this analysis is to establish the relationship between the microstructure and toughness of some experimental results obtained over the past decades for weld metals with tensile strength from 400 to 1,000 MPa. The analysis was carried out using the methodology proposed in [32], to verify its eff ectiveness and explanation impact toughness behavior. Research results of several authors and discussion Infl uence of carbon equivalent on tensile strength and impact toughness of weld metals Figure 1 shows the eff ect of carbon equivalent on the strength and toughness of the weld metal from the review work [32]. It was shown in [32] that Ceq has a high dependency on the tensile strength of the weld metals (fi gure 1, a), and some studies have shown an almost linear increase in the ultimate strength of the weld metal with increasing Ceq. It can be seen that with an increase in the strength of the metal, a high scatter of values is observed, which may be due to diff erent cooling rates, since the high hardenability of the alloys contributes to the same microstructure of the entire weld metal. However, small deviations in cooling rates cause signifi cant
OBRABOTKAMETALLOV MATERIAL SCIENCE Vol. 26 No. 1 2024 changes in the amount of martensite, bainite and acicular ferrite [30]. Figure 2 shows that a high scatter band is observed when high-strength weld metals are cooled for diff erent times in the temperature range 800–500 °C [4]. Standards [5–10] allow a wider range of alloying and micro-alloying elements, and therefore each manufacturer off ers its own chemistry to achieve qualifi cation requirements. Carbon equivalent (Ceq) was included in the standard [5] because it is generally related to hardenability. Limits for Ceq were calculated based on the minimum and maximum alloying element contents. Therefore, a lower Ceq value is always preferable, indicating good weldability. The American Petroleum Institute has adopted two formulas (CEIIW and CE Pcm) [5] to determine the carbon equivalent limit for API PSL 2 pipe steel. The CEIIW formula is provided by the International Welding Institute and is commonly used for carbon and carbon-manganese steels. In Europe Pcm is the critical parameter of the metal. CE Pcm is taken from the documents of the Japan Society of Welding Engineers. CE Pcm was proposed specifi cally for testing the weldability of high-strength steels. The balance of superior strength and toughness can be disrupted following thermal cycling that occurs during welding, causing poor toughness in the heat-aff ected zone (HAZ) [11–19]. General welding issues Modern steels with high strength and high impact toughness are widely used in pipelines, shipbuilding and various manufacturing industries [2, 3]. Changes in steel production technology and the steel rolling process pose a challenge to the production of welding consumables and joining technology. It is important to note that, in contrast to the production of wrought steel, the strength and toughness of weld metals, as a rule, should be achieved through alloying [2–4]. As a consequence, due to the complexity of welding processes and the limitation of heat input and, consequently, cooling rates, the toughness of the weld metal at low temperatures is lower than that of the base metal [3, 4]. In addition [2–4], the microstructure of weld metals with a yield strength of 600 MPa and above consists mainly of bainite and martensite, rather than predominantly acicular ferrite. Therefore, the calculation of the basic composition of the weld metal should a b Fig. 1. The eff ect of the carbon equivalent on the ultimate tensile strength (a) and impact strength at 20 °C of weld metals (b) [32] Fig. 2. The eff ect of weld metal cooling rate (Δt8/5) on the ultimate tensile strength of high-strength pipeline steels [4]
OBRABOTKAMETALLOV MATERIAL SCIENCE Vol. 26 No. 1 2024 be diff erent for each case [2]. In fact, for those applications where the strength of the weld metal consisting of acicular ferrite is insuffi cient, it is necessary to add special strengthening elements for solid solution and other alloying elements to retard the austenite/ferrite transformation and produce martensitic welds with the required high strength. In work [4], specimens obtained by the SMAW (Submerged Metal Arc Welding) method were studied; automatic submerged metal arc welding and GMAW (Gas Metal Arc Welding) is a designation used to indicate the use of the MIG/MAG method in automatic (robotic) welding. The authors wanted to evaluate (fi gure 3) whether the use of the GMAW process could improve the weld performance of high-strength steels while maintaining good quality even at lower levels of reheat. It was found that a good relationship between mechanical strength and toughness could be obtained. Multi-pass welding is widely used in pipe manufacturing, circular welding of pipe butt joints, and in-service welding. For automatic welding of large diameter pipes, the root welding method by an internal welding machine and the cap welding of external welding machine are usually used [18]. The pipe neck group is fi rst welded to the inner root of the pipeline using a welding robot, and then the weld root is welded (hot pass), then the fi ller and lining layers of the joint are welded, as shown in fi gure 4, where layer 0 is a root weld, layers 1–6 are fi lling ones, and layers 7–8 are cap ones. The fi rst layer completes the fusion of the root weld. Due to the large number of fi lling layers, the likelihood of defects greatly increases. Thermal cycles encountered during welding are characterized by a range of peak temperatures that can alter the microstructure and properties of the HAZ compared to the base metal. It was found that the supercritical (reheated above Ac3) and subcritical (reheated below Ac1) regions resulting from the second thermal cycle retain toughness properties comparable to the original ones. Among all the sub-zones of the HAZ in multi-pass welding, the IC-CGHAZ (i.e., the pre-existing CGHAZ reheated to the temperature range between Ac1 and Ac3 in a subsequent weld) is considered to be subject to the most signifi cant degradation in toughness [11–18]. This is confi rmed by the work of the authors [29, Fig. 3. The relationship between the mechanical strength and impact toughness of the weld deposit of high strength steels in comparison with several works by SMAW (Submerged Metal Arc Welding) and GMAW (Gas Metal Arc Welding) method [4] Fig. 4. Transverse microsection of an annular welded joint of pipes with narrow edge cutting [18]
OBRABOTKAMETALLOV MATERIAL SCIENCE Vol. 26 No. 1 2024 30], and the signifi cant deterioration of IC-CGHAZ is associated with the presence of a blocky component of the martensite-austenite grain boundary (M-A). It is well known that the HAZ is the weakest part of the welded joint and determines the safety of pipeline operation. In particular, the lowest impact toughness was obtained in the coarse-grained microstructure of the HAZ [16–19], which is adjacent to the fusion line of the weld. Low impact toughness of the HAZ at low temperatures is the main problem limiting the use of highquality steels for pipelines [2–4, 11–29]. Description of microstructures It should be recognized that the problems of welding high-strength steels are far from being solved. For example, it is known that non-metallic inclusions formed in the cast weld metal have two opposite eff ects on impact strength [30]. Firstly, inclusions act as initiation sites for both ductile and shear fractures [29–31], and secondly, it can promote the formation of acicular ferrite, which is recognized as the most optimal microstructure [31–42]. One of the basic requirements for pipeline joints is to obtain weld metal of equal or higher strength than the base material to avoid localized deformation or failure of the weld under load. However, suffi cient strength is also required, which is usually verifi ed using Charpy impact tests. A common solution is to develop the weld metal to produce acicular ferrite (AF) in the metal structure, which provides a balance between strength and toughness [28, 29]. This fact has stimulated extensive research into the mechanisms of AF formation in weld metals and the determination of what factors control its formation [11–39]. The key factor in the formation of AF is the chemical composition of the consumable welding wire, both from the point of view of the isolated eff ect of each element and the combined eff ect of the general composition [29–38]. In a review [32] on the formation of acicular ferrite in carbon-manganese surfacing, it was reported that the following elements infl uence the formation of acicular ferrite: C, Mn, Si, Ni, Ti, Al, Mo and Nb. The eff ect of austenitization temperature in the range of 850–1,000 °C on acicular ferritic transformation in Cr65 HSLA pipeline steel was investigated [31]. As shown in fi gure 5, the initial and fi nal phase transformation temperatures during continuous cooling, namely Ar1 and Ar3, respectively, decreased with increasing austenitization temperature. This result suggests that increasing austenitization improves the stability of austenite during cooling and thus delays the decomposition of austenite. The decomposition products of austenite in Cr65 steel consist mainly of polygonal ferrite, pearlite, acicular ferrite, etc. [30]. Increasing the austenitization temperature promotes the formation of acicular ferrite and prevents the formation of pearlite and polygonal ferrite (see fi gure 6). A higher austenitization temperature leads to more suffi cient dissolution of carbide-forming elements such as Nb, V and Ti, as well as more suffi cient homogenization in the austenite [30]. The authors of [31] believe that dissolved alloy elements improve the stability of metastable austenite. Thus, according to the authors, the decomposition of austenite is delayed to a lower temperature, which is also confi rmed by fi g. 3. As a diff usionless reaction [32], the acicular ferrite transformation is more likely to occur at a relatively low temperature than the diff usion-controlled pearlite or polygonal ferrite transformation [33], since the rate of atomic diff usion decreases with decreasing temperature. Fig. 5. The fractions of phase transformation determined by dilatometric measurements as a function of temperature during continuous cooling in X65 pipeline steel specimens austenitized at diff erent temperatures from 850 °C to 1,000 °C [30]
OBRABOTKAMETALLOV MATERIAL SCIENCE Vol. 26 No. 1 2024 Copper has been reported to contribute to the formation of AF when using manual arc welding [32]. Many elements will combine with the oxygen present in the weld metal, which can be controlled by shielding gas and/or the composition of the weld metal. The oxygen reaction infl uences AF formation by either promoting or inhibiting the formation of nonmetallic inclusions such as oxides. Some authors argue [29–48] that oxides act as nucleation sites for AF, so an increase in oxygen content favors the formation of AF. For example, it was reported [33, 34] that increasing oxygen content to 300 ppm changed the weld metal of Widmanstätten side plates to AF microstructure [33, 34]. The formation of AF is also facilitated by coarse austenite grains with a large number of inclusions with a diameter of more than 0.2 μm. The details of the formation of AF are now well described as a variant of the bainite structure in numerous works by Bhadeshia and his students [48], where it is shown that it is a special variant that depends on intragranular formation [33]. Thus, it is necessary to achieve suffi cient preliminary austenite grain size and number density of non-metallic inclusions of favorable chemical composition, especially based on titanium oxides. However, it has also been noted in many works that if the amount of non-metallic inclusions reaches a certain level depending on the oxygen content, it has a detrimental eff ect on the toughness as the crack initiation sites outweigh the benefi ts of achieving a fi ne AF structure. Two welding materials suitable for joining Cr80 steel pipes are compared in terms of weld metal microstructure, hardness, toughness and tensile properties [35]. The chemical composition of the consumables was similar: one of the consumables had a rich chemical composition of the wire and contained higher alloying additives C, Ni, Ti compared to the depleted wire. Deposit welding was performed using a gas arc welding (GMAW) process system to achieve the same heat input of 0.66 kJ/mm. The results showed that for both wires, the microstructure of the weld metal was mainly composed of acicular ferrite. a b c d Fig. 6. Optical micrographs of X65 pipeline steel specimens after continuous cooling with diff erent austenitization temperatures: 850 °C (a), 900 °C (b), 950 °C (c) and 1,000 °C (d) [30]
OBRABOTKAMETALLOV MATERIAL SCIENCE Vol. 26 No. 1 2024 Consumables with a richer chemical composition (C, Ni and Ti) showed higher strength and hardness due to the fi ner microstructure of the fi nal weld metal; however, Charpy impact test results showed that the depleted chemical wire had higher impact strength at low temperature. Since both weld metals had a similar acicular ferrite structure, the lower toughness of the richer weld was attributed to the presence of titanium inclusions, which could become crack initiation sites. The eff ect of welding method and preheating on the weld metals of pipeline steels was investigated in the work of HSLA steel. It is known that post-welding treatment reduces the strength characteristics of the metal of the welded joint [2, 4]. The results of studies [4] assessing the preheating up to 200 °C and welding treatment of welded joints showed a tendency towards a decrease in mechanical strength and an increase in impact strength as a consequence of some important aspects, such as a lower percentage of martensite, coarsening of the microstructure and a higher proportion of high-angle boundaries (> 15 %). Longer cooling times (time spent in the temperature range of 800–500 °C) show a tendency to improve the impact toughness and reduce the mechanical strength of deposited metals in high-strength steels. Microstructure features aff ecting the impact strength of weld metals Multi-pass welding is required to connect the main pipes. This leads to overheating of the HAZ. This creates its own peculiarities of thermal eff ects on the metal and, as a result, non-classical phase and structural transformations with sharp temperature and stress gradients. The HAZ zone can be divided into coarse-grained HAZ (CGHAZ), fi ne-grained HAZ (FGHAZ), intercritical HAZ (ICHAZ) and subcritical HAZ (SCHAZ), when a single thermal cycle is applied to weld the material [43]. When a second weld bead is applied over an existing one, it results in the formation of a plurality of reheated HAZ structures that are characterized by corresponding second peak temperatures and include supercritical, intercritical and subcritical structures. The strength and toughness of HSLA pipeline steel can degrade signifi cantly after one or two welding thermal cycles, so CGHAZ with intercritical reheating (ICR) are often considered to be the weakest link or most fragile area of the weld joint. A schematic representation of a weld with diff erent heat-aff ected zones is shown in fi gure 7 [38]. Various metallurgical factors such as austenite grain size and bainite stack size, as well as the size, shape and distribution of any second phase (carbide or martensitic-austenitic) can infl uence fracture toughness. In particular, the presence of so-called martensitic-austenitic (MA) constituents formed in ICR CGHAZ plays a decisive role in the fracture toughness at low temperatures. Fig. 7. Schematic representation of microstructures in the heat-aff ected zone of multi-pass welds [43]
OBRABOTKAMETALLOV MATERIAL SCIENCE Vol. 26 No. 1 2024 Although MA has been widely studied in recent decades, the eff ect of cooling rate on its volume fraction remains controversial [23–29]. Some researchers have shown that increasing the cooling rate increases the MA fraction, while others have shown, on the contrary, [32–35] that a slower cooling rate decreases the MA fraction. Works [38–48] showed that for various steels there is an increase in the fraction of MA at a lower cooling rate. In addition to the eff ect of cooling rate on MA fraction, the eff ect of MA grain size, morphology and distribution on impact toughness has also not been established. This is largely due to the complex factors that determine impact strength, including the fraction, size, substructure and morphology of MA. It is generally accepted that MA reduces the impact strength of pipeline steel [4]. A slower cooling rate results in a coarser MA structure, resulting in poor toughness properties. In the work [43] it is reported that the formation of lath type (thin MA), associated with poor toughness, occurs at slower cooling rates, while block MA is formed at higher cooling rates. It is important for further analysis to interpret the microstructure of steel after welding because this is controversial because constituents that are part of the same primary structure may appear morphologically diff erent depending on the viewing plane (fi gs. 8, 9), and some structures may have similar morphological features but present diff erent mechanical properties [44–46]. Fig. 7 shows an overview of the evolution of the major components present in the weld metal as observed by optical microscopy (OM). This fi gure shows that the microstructure changes continuously with increasing carbon equivalent (Ceq). A mixture of acicular ferrite (AF), primary ferrite (PF) and ferrite with a second phase (FS) in the columnar region has lower strength values. In contrast, the reheat region is dominated by polygonal ferrite. In addition to the tendency to have a mixture of martensite and bainite with a higher content of alloying elements due to increased hardenability, it is worth noting the presence of similar components for both columnar and reheated areas. The terminology of microstructural constituents observed in weld metals has been very confusing [35], with diff erent terms being used to refer to the same constituent. This lack of clarity prompted the International Institute of Welding (IIW) to develop a general framework for microstructure quantifi cation [36] in the 1980s, where components were easily identifi ed using optical microscopy (OM). Another critical issue relates to the low resolution of optical microscopy for refi ning the constituents of refi ned weld metals, even when using higher magnifi cation than recommended by the IIW [38, 39]. To solve this problem, scanning electron microscopy (SEM) has been widely used in recent decades, mainly to separate bainite and martensite and evaluate micro-phases. However, sometimes even this method has limitations in distinguishing the overall microstructure. This occurs mainly for weld metals with a tensile strength greater than 600 MPa, where a mixed microstructure consisting of acicular ferrite, bainite (ferrite with a second phase) and martensite predominates. To ensure proper resolution in the study of microstructure, the EBSD technique is used as an additional tool [11, 12, 32]. This method has been considered as an interesting alternative [32–40] to overcome the shortcomings of optical microscopy. This method, which provides valuable grain boundary information, is useful for refi ned microstructures to confi rm constituents such as acicular ferrite, bainite and martensite. The high clarity provided by EBSD, especially at grain boundaries, is useful for separating acicular ferrite and bainite (second-phase ferrite). Regarding the assessment of MA components and inclusions, SEM analysis is more suitable for this task [36–39]. Thus, it is believed [11, 12, 24, 25, 32–40] that a combination of OM, SEM and EBSD methods provides the best methodology for the study of metal welds in C-Mn steel in the presence of a refi ned microstructure. In [38], the author examined the IIW scheme for the main structures that develop during reduction and shear transformation in steels. However, he noted that questions remain to be resolved regarding reaction kinetics, especially elucidating the growth mechanisms of bainite, which could lead to greater precision in distinguishing bainite from other phases. In [38], a critical review is presented to clarify the existing confusion in the literature regarding bainite and acicular ferrite due to the similarity in appearance of these two microstructural constituents observed under an optical microscope. The works [44-48] present a description of the microstructural components in relation to low-carbon pipe steels.
OBRABOTKAMETALLOV MATERIAL SCIENCE Vol. 26 No. 1 2024 Fig. 8. Microstructures observed in weld metals with increase in tensile strength after etching with Nital 2 %. Magnifi cation: 1,000× (OM). Where: AF – acicular ferrite; PF – primary ferrite; FS – second phase ferrite; M – martensite [32]
OBRABOTKAMETALLOV MATERIAL SCIENCE Vol. 26 No. 1 2024 Fig. 9. SEM images of weld metals are shown after etching with Nital 2 %. Where: AF – acircular ferrite; PF – primary ferrite; FS – ferrite with second phase; M – martensite [32]
OBRABOTKAMETALLOV MATERIAL SCIENCE Vol. 26 No. 1 2024 Our analysis shows that, in relation to low-carbon pipe steels, the weld metal can have the following microstructures [11–49]: – primary ferrite, which is nucleated at the boundaries of the initial austenite grains (allotriomorphic ferrite) and to a lesser extent inside the austenite grains (euhedral ferrite), where non-metallic inclusions (NI) are presented [39–42]. Primary ferrite with nuclei at grain boundaries is formed during cooling in the temperature range of 1,000 and 650 °C [20–34]; – side plates of ferrite [34, 39–42] (separated by low-angle boundaries) are formed during cooling at temperatures from 750 to 650 °C, also at the boundaries of primary austenite grains [29]; – acicular ferrite [34, 39–42] is heterogeneously nucleated on the surface of non-metallic inclusions during the austenite-ferrite transition. As the transformation proceeds, ferrite grains diverge in diff erent directions, creating a chaotic structure [29, 30] of crystallographically misoriented plates approximately 5–15 μm long and 1–3 μm wide [17–29, 39–42]. The temperature range, over which acicular ferrite is formed, depends on the overall composition and cooling rate across the transformation temperature range, but is typically in the range of 750–560 °C [34, 35]. – bainite grows in the form of individual plates or subunits [48], which can form bundles of parallel ferrite laths [34]. It can be classifi ed as upper or lower bainite depending on the transformation temperature. In upper bainite, carbon is deposited as cementite (Fe3C) between bainitic ferrite plates (bundles) [48]. In lower bainite, the ferrite becomes supersaturated with carbon, and some carbide precipitation occurs within the ferrite subunits as well as between it [43]. The initial temperature of bainite depends on the composition and cooling rate, but is usually on the order of 560 °C [48–67]. The nucleation effi ciency of nonmetallic inclusions in modern low-alloy steel weld metals is such that the colony size of intragranular bainite is similar to the size of acicular ferrite in C-Mn steel weld metal [29]. Consequently, when examined under an optical microscope, the colonies within granular bainite are very similar to the type of acicular ferrite with which it is confused in the literature [41–44]. Some authors [44, 45] use the term granular bainite, which does not diff er from lath bainite in terms of the transformation mechanism, although granular bainite packets form at relatively higher temperatures and mainly consist of wide parallel laths, while lath bainite packets form at relatively lower temperatures and consist of thin parallel slats; – transformation of pearlite can occur at the boundaries of austenite grains or in such inhomogeneities as inclusions. At high transformation temperatures, pearlite forms nodules of alternating plates of ferrite and cementite, which can be quite large. As the transformation temperature decreases, the pearlite sheets become increasingly thinner until the structure becomes indistinguishable under a light microscope. Alternatively, distorted pearlite plates may appear as a virtually indistinguishable ferrite/carbide aggregate [53–56]. Lamellar pearlite, FC(P) in the IIW classifi cation scheme [35], can be confused with martensite if the ferrite/cement laminae are indistinguishable under the optical microscope [2, 41–44]. – martensite is formed as a result of a rapid and diff usion-free transformation, in which carbon remains in solution [43]. Martensite can occur in the form of laths or plates. The substructure of lath martensite is characterized by a high density of dislocations located in cells, where each martensite plate consists of many dislocation cells. The substructure of lamellar martensite consists of very small twins, i.e. twinned martensite [42–45]. The mechanisms of formation of the components are not discussed in this work, since there is extensive material on this topic in the literature [33–67]. It is noted in [67] that, unlike metals of single-pass welds, metals of multi-pass welds contain in each bead (except for the last bead) a large proportion of overheated areas, which, due to subsequent beads, are reheated to a temperature above Ac3. The eff ect of multiple welding passes on C-Mn and low-alloy steel deposits is very complex since the proportion of columnar and recrystallized regions and its corresponding microstructures depend on various parameters such as heat input, temperature between passes and chemical composition [29]. The previous columnar morphology changes during the reheating process, resulting in a heterogeneous microstructure that aff ects the performance of the welded joint [4, 29, 32].
OBRABOTKAMETALLOV MATERIAL SCIENCE Vol. 26 No. 1 2024 Mechanical properties The authors [4] stated that only a few studies have examined the mechanical properties of reheated weld metals. The results are still inconsistent as it depends on several factors such as the amount of acicular ferrite and the presence of MA components. The authors of [29–33] noted that understanding the variation in toughness in the multi-pass weld metal of C-Mn steels is very diffi cult, even if the eff ect of reheating due to multiple passes is taken into account. Similarly, the authors of [29, 38–42] suggested that signifi cant changes in the toughness of C-Mn weld metals are due to the microstructural features existing in the Charpy-V notch, which are the combined result of the chemical composition, welding procedure, deposition sequence and specifi c welding methods. In addition to the factors mentioned above, it is critical to consider the position of the Charpy-V notch in relation to the proportion of weld metal reheated. A specifi c assessment should be made for each case. The author [48] observed that although complete recrystallization was observed for two intersecting regions per layer and the proportion of reheated regions was about 75–80 % for three layers, the same toughness was obtained for both sequences, depending on the Mn content. In general, toughness increases when the fraction of recrystallized region increases due to the predominance of polygonal ferrite, microstructure refi nement, or tempering eff ects during subsequent deposition [41–58, 67]. However, some data suggests that this property deteriorates with extensive segregation [29, 30] or the presence of micro-phases located along the grain boundaries of the previous austenite [47, 48]. Another negative contribution is associated with a decrease in the proportion of acicular ferrite due to the smaller size of the precursor equiaxed austenite grains in the reheated weld metal [46–53]. Figure 10 shows an OM image of a Charpy-V notch, where the ratio of columnar and reheated regions can be easily determined for C-Mn weld metals since these regions are well defi ned. For more alloyed weld metals, this distinction can be more complex. In this case, several polishing and etching steps may be required to enhance the contrast between the areas. The authors of [54] noted that only a few studies have focused on the microstructure and toughness of actual weld metal. This is because it is very diffi cult to analyze its correlation using a real weldment, and the precise determination of the correlation between the ring-type MA component in the weld metal reheat zone and the toughness is still uncertain. At the same time, after metallographic studies, when an accurate characterization of the microstructure has been obtained, it is possible to assess impact strength based on the following criteria. (1) Reheat. This criterion is less representative in many of the studies analyzed because the same proportion of recrystallization was obtained for all deposits; (2) Microstructure. EBSD results confi rm this trend, showing that fi ner microstructure has a higher frequency of high-angle boundaries (HABs), which can eff ectively cause the propagation of cleavage cracks to defl ect or stop [32–46]. The same behavior was noted for the reheating region of the refi ned grains, where polygonal ferrite predominates; (3) No metallic inclusions. It is known that non-metallic inclusions can have two opposite eff ects on impact toughness [11, 12]. One of it is that inclusions act as crack initiation sites, both plastic and cleavage. Secondly, it can promote the formation of acicular ferrite. It has been observed that increasing the Ti content promotes the formation of inclusions suffi cient to support the formation of purifi ed acicular ferrite, in agreement with other works [3, 11, 12, 32, 36]. Fig. 10. Optical microscopy at low magnifi cation of the Charpy-V notch position for C-Mn weld metal after etching with Nital 2 % [32]
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