High alloy steel: there are alloying elements of more than 8% by weight of total other than carbon and iron is classified as high alloy steel.
From: Civil Engineering Materials , 2021
History and organization of codes
Maurice Stewart , in Surface Production Operations, 2021
2.12.4.3.3 High-alloy steels
High-alloy steels are usually called stainless steels. They are known for high levels of corrosion resistance. Stainless steels used for pressure vessels and piping construction are divided into three groups:
•
Austenitic stainless steels
•
Ferritic stainless steels
•
Martensitic stainless steels
Steels from these three groups have the following characteristics.
2.12.4.3.3.1 Austenitic stainless steels
Austenitic stainless steels consist of chromium-nickel (300 series) and chromium-nickel-manganese (200 series). They are nonmagnetic, highly corrosive resistant, and hardenable only by cold working (strength and hardness can be increased at the expense of ductility). They are the only steels assigned allowable stresses in the code for temperatures above 1200°F (649°C), up to 1500°F (816°C). The most commonly used are Grades 304 and 316. For welded construction, low-carbon Grades 304L and 316L or stabilized Grades 321 and 347 are usually specified. Higher chromium content steels, Grades 309 and 310, are resistant to oxidation and sulfur attack up to 2000°F (1093°C).
2.12.4.3.3.2 Ferritic stainless steels
Ferritic stainless steels are straight chromium stainless steels with a minimum of 10% chromium. They are nonhardenable by heat treatment. They are seldom used in vessel construction except for corrosion-resistant lining or cladding. They are typically used for internal trays for less-corrosive environments. Common grades are Grades 405 and 430.
2.12.4.3.3.3 Martensitic stainless steels
Martensitic stainless steels contain 11% to 16% chromium with sufficient carbon to be hardenable (less than 1%). They are hardenable by heat treatment. They are (1) the least corrosion resistant of the stainless steels. They are difficult to form and require heat treatment after welding. They are rarely used as construction materials for pressure vessel parts, with the exception of internal linings and trays, because of poor weldability.
P.D. Pattillo , in Elements of Oil and Gas Well Tubular Design, 2018
Example problem—kinematic hardening with von Mises yield criterion
When manufacturing tubing of some high alloy steels, the work hardening accompanying forming of the tube can result in a product that has a tensile axial yield stress that is higher than its compressive axial yield stress. Consider the two-dimensional representations of the three-dimensional von Mises yield surface that have as their abscissa either the effective stress (Section 6.3.4.1, Fig. 6.11) or the sum of the axial stress and internal pressure (Section 6.3.4.2, Fig. 6.13). We wish to determine how these yield surfaces should be altered to model anisotropic yield in the axial direction only. We assume the tensile yield stress is
and the compressive yield stress is .
One heuristic method of altering the yield surfaces is fairly accurate and involves some variation of the following technique:
•
Begin with a yield surface for an isotropic material with
.
•
Locate a new point on the abscissa to represent the compressive yield stress
.
•
For the entire compressive side of the yield envelope, proportionately alter each ordinate value by the formula
(6.131)
For example, the new y value for a compressive stress that is one-half of compressive yield in the current problem would be
(6.132)
that is, half way to compressive yield on the altered yield surface equals half way to compressive yield on the original yield surface. We have essentially pushed the isotropic compressive yield value to the anisotropic yield value, allowing the yield surface to shift proportionately for all negative values on the abscissa. This method is illustrated in Fig. 6.17 for the two yield expressions from Sections 6.3.4.1 and 6.3.4.2.
Figure 6.17. Alteration of von Mises yield surface to accommodate axial anisotropic yield using a heuristic method. Positive abscissa values of the yield surface are unaltered. During manufacture the pressures vanish so both abscissas become Σzz. (A) Effective stress; (B) Axial stress +pi.
Let us, as an alternative to the heuristic model, seek a yield surface which results from the application of kinematic hardening. We could repeat the calculation for plastic strain as we did in the isotropic hardening example of Section 6.4.3.1. Our primary concern, however, is the translation of the yield surface, so we will limit our consideration to that aspect of the plastic deformation.
Initially, in Eq. (6.130). Further, since during manufacturing there is no internal or external pressure, the abscissa of both of the presentations considered here is . This means that in both presentations the drawing process takes place along the abscissa and the only nonzero component of translation is . That is, even it we do not evaluate , we know that in each case the translation will occur along the -axis. Further, we know that (a) the shape and size of the yield surface do not change and (b) the difference between the yield in tension and compression is
. Assuming we started the drawing process with an isotropic tube, this implies the initial yield stress—both in tension and compression—was . The drawing process has translated the sample along the positive -axis.
Fig. 6.18 compares the results of kinematic hardening with the heuristic model just presented. In the effective stress formulation of the initial yield surface shown on the left-hand diagram in the figure, the difference is small. The right-hand diagram indicates a greater difference, particularly in the vicinity of the ordinate boundary between quadrants 1 and 2 of the graph. Both diagrams illustrate the size-preserving translation the yield surface undergoes under kinematic hardening beyond initial yield.
Figure 6.18. Alteration of von Mises yield surface to accommodate axial anisotropic yield using kinematic hardening. During manufacture the pressures vanish so both abscissas become Σzz. (A) Effective stress; (B) Axial stress +pi.
J.M. Antonini , in Comprehensive Materials Processing, 2014
8.04.3.1.2.2 Chromium
The welding of stainless steel and high alloy steels produce fumes that contain chromium. Chromium in welding fumes can exist in different oxidation states such as Cr6+ and trivalent chromium (Cr3+). Cr3+ has been considered to be of a low-order toxicity because it does not enter the cells of the body, whereas Cr6+ can enter cells and has been found to be quite toxic (29,30). Significant quantities of Cr6+ have been measured in fumes during stainless steel welding. The selection of the type of shielding gas and weld process can influence the concentration of Cr6+ produced (24). Cr6+ has been classified as a human carcinogen (31), and welding fumes that contain chromium have been shown to be mutagenic in cell-based studies (32,33). Epidemiology studies have indicated a possible increase in mortality from lung cancer among stainless steel welders (34–36). Because of the carcinogenic potential of stainless steel welding fumes, federal legislation in 2006 reduced the permissible workplace exposure limit (PEL) for chromium by an order of magnitude from 52 to 5 μg m–3(37). The decision to lower the exposure limit was based on the finding that employees exposed to Cr6+ faced an increased risk for significant health effects.
K. Aasberg-Petersen , ... A.P. Steynberg , in Studies in Surface Science and Catalysis, 2004
Tube Design
A steam reformer may contain up to about 1000 high alloy steel tubes filled with catalyst. The outer diameter is typically 100 – 180 mm, the tube wall thickness is 8 – 20 mm, and the heated length may be 10 – 14 m. The traditional materials used for reformer tubes were HK 40 (25 Cr 20 Ni) or IN 519 (24 Cr 24 Ni Nb) [29, 43, 44]. However, the demand for better creep rupture properties led to the development of new alloys with greater strength. The development in the strength of reformer tube materials is illustrated in Fig. 20 [44].
Figure 20. Creep rupture strength for different reformer tube materials [44]
Today, the preferred tube materials are the so-called 'microalloys', typically 25 Cr 35 Ni Nb Ti [43]. The other types of materials will only be used in cases where the strength of the microalloys cannot be used fully, as there is a lower limit for the wall thickness of the tubes.
Reformer tubes are normally designed according to the 'remaining life assessment' technique, API-530 [45], for an average lifetime before creep rupture of 100,000 hours. The main parameters in the design are the design pressure; the design temperature and the creep rupture strength of the material used. However, the determination of these parameters can be ambiguous, and each reforming technology licensor applies in-house procedures to determine the parameters and to introduce necessary design margins.
The calculation of the design temperature is demanding since it requires detailed understanding of the heat transfer. This includes heat transfer to the individual tube by radiation from the furnace internals including furnace walls and neighbouring tubes as well as heat transfer by convection from gas to tube wall, by conduction through the tube wall, and by convection from the inner tube wall to the catalyst and the reacting gas. Secondly, an understanding of reaction kinetics, catalyst ageing, and heat and mass transfer (radial and axial) in the catalyst bed, etc. is required. The interplay between catalyst, reacting gas and reformer tube is also essential for the prediction of the limits for carbon formation.
Yusheng Shi , ... Ying Chen , in Materials for Additive Manufacturing, 2021
5.6.2.2 Heat treatment process and properties
The deposited metal of the flux-cored wire MDY501 is high-alloy steel with high strength and low toughness. The heat treatment method further enhanced the performance of the formed components, mainly improved the toughness. Meanwhile, the heating temperature caused the secondary hardening of carbides of V and Cr, so that the toughness could be greatly increased without sacrificing too much hardness, thereby satisfying the usage requirements.
Table 5.31 shows the Rockwell hardness values obtained after heat treatment of the post-deposited metal at a series temperature of 480°C, 500°C, 520°C, and 560°C. It can be seen from the table that the average hardness of self-designed high-alloy steel flux-cored wire MDY501 depositing layer was 49.1 HRC, and its hardness reached the highest value after heat treated at 500°C, which was 45.8 HRC. The hardness value of the deposited metal decreased when heat treatment temperature increased further because of the atom migration and diffusion during the recovery process, and the dislocation density decreased, the number of internal defects of the metal and the distortion energy decreased, and the internal stress was eliminated, thereby causing a decrease in hardness.
Table 5.31. Hardness values of as-deposited and heat treated of high-alloy steel samples (HRC).
Testing sample no.
1
2
3
4
5
Average value
1
As-deposited480°C heat treated
49.6
48.6
48.5
48.3
50
49
43
41.2
43
41.5
43.2
42.4
2
As-deposited500°C heat treated
48
48.6
48.4
47.9
49.1
48.4
45.5
46
44.5
47.2
46
45.8
3
As-deposited520°C heat treated
49.9
50.5
49.4
49.8
49.5
49.82
43
41.5
43.5
41
42
42.2
4
As-deposited560°C heat treated
49.2
49.1
49.1
49
49.4
49.16
39.1
39
38.9
38
38.5
38.7
The effect of heat treatment temperature on the hardness of the depositing layer is shown in Fig. 5.165. It can be seen from the figure that when the heat treatment temperature was in the range of 480°C–500°C, the hardness of the deposited metal showed an upward trend, mainly due to the precipitation of carbides in the martensite phase, increasing the non-uniform nucleation rate, refining the grains, and making the lath martensite transform into ferrite equiaxed grains. The hardness of the deposited metal reached a peak at 500°C, which was 45.8 HRC, indicating that the deposited metal had secondary tempering and hardening when heat treatment was performed under this temperature specification; when the heat treatment temperature was greater than 500°C, the hardness of the welding layer began to decrease, and the hardness value decreased to 38.7 HRC at 560°C. The hardness of the deposited metal was mainly determined by the carbon content in the martensite. When the temperature was greater than 560°C, the martensite decomposed and precipitated carbides, resulting in a continuous decrease in carbon content.
Figure 5.165. Effects of heat treatment temperature on hardness of depositing layer.
Fig. 5.166 shows the metallographic structure of the deposited metal after heat treatment at different temperatures (480°C, 500°C, 520°C and 560°C), respectively.
Figure 5.166. Metallographic structure of deposited metal after different heat treatment temperatures.
As shown in Fig. 5.167, when the heat treatment temperature was 480°C, the microstructure of the deposited metal was mainly tempered martensite and a small amount of retained austenite, the martensite was lath shaped and the carbide was an alloy carbide formed by Cr, V and C. At 500°C, the precipitation of carbides increased, and the particles were small and dispersedly distributed. As the heat treatment temperature rise to 520°C, the martensite in the welding layer metal decomposed, precipitated carbides and diffused along the grain boundaries, thus promoting nucleation and inhibiting grain nucleation, and the lath martensite transformed to ferrite equiaxed grains; when the heat treatment temperature reached 560°C, the cementite aggregated and the ferrite phase appeared recovery and recrystallization, but due to the decomposition of martensite, a large amount of carbide dispersed phase was precipitated, so that the ferrite in the metal of welding layer could not be sufficiently recrystallized to form large number of ferrite equiaxed grains.
Figure 5.167. ESEM scan of WAAM samples heat treated at (A) 480°C, (B) 500°C and (C) 520°C.
It can be seen from the ESEM photos that when the heat treatment temperature was 500°C, the martensite was thin and uniform, the amount of the second phase particle was the largest, the particles were fine and finely distributed; lath martensite increased no matter whether the temperature was high or low, and the performance was not optimal. Through this series of comparative tests, the optimum heat treatment temperature of the flux-core wire can be obtained at about 500°C. In this temperature range, the hardness would reach a peak value due to "secondary hardening" effect, and the toughness was also enhanced, so that optimized microstructure and performance were obtained.
Morphologies and mechanisms of metal dusting on high-alloy steels
P SZAKÁLOS , in Corrosion by Carbon and Nitrogen, 2007
Publisher Summary
This chapter describes the morphologies and mechanisms of Metal Dusting (MD) on high-alloy steels. It concludes that in contrast to pure iron, it is not possible to explain the MD process on high-alloyed steels by a solely carbon-induced corrosion. The dissolution of the alloyed carbides by selective oxidation during metal dusting is one of the key mechanisms, denoted Type III. The driving force for MD in pure iron and low-alloyed steels is explained by separating the initial, cementite-formation stages (Type I), from the steady state MD process, which may be described as disintegration of supersaturated ferrite by graphite formation, denoted the Type II mechanism. The important conclusion is that there is no inexplicable step or phenomenon included in the different metal dusting processes that occurs in pure metals or engineering alloys. In fact, all the observed and reported observations are readily explained with four fundamental mechanisms, Types I, II, III, and IV, which follows the known thermodynamics in the respective system Me–C–H–O. The Type IV mechanism may be described as continued disintegration of the corrosion products resulting in nanoparticles and carbon nanotube formation.
α-Fe layer formation during metal dusting of iron in CO–H2–H2O gas mixtures
J ZHANG , ... G INDEN , in Corrosion by Carbon and Nitrogen, 2007
11.1 Introduction
The catastrophic carburisation process, metal dusting, attacks iron, low- and high-alloy steels, and Co or Ni-based alloys in strongly carburising gas atmospheres with carbon activity aC > 1 at elevated temperatures. The processes and mechanisms of metal dusting have been studied over decades, and substantial, but not complete, understanding has been achieved. A general mechanism [1–6] proposed for iron-based alloys involves a supersaturation of iron with carbon and subsequent growth of a cementite layer at the surface which acts as a barrier for further carbon transfer. Consequently, graphite is deposited on the cementite surface, lowering the carbon activity to one in the Fe3C/graphite interface and initiating cementite decomposition. The particles formed during the decomposition process then strongly catalyse further carbon deposition causing a vast growth of the reaction product 'coke'.
The formation of iron carbide and its decomposition have been reported extensively in metal dusting of iron and in cementite production in the iron-making process [1–8]. It was found that the rates of cementite formation and its decomposition depend on the composition of the reaction gases and on temperature [6–8]. Recently the formation of iron particles or even an iron layer between cementite and graphite was observed on iron at 700 °C in H2–CO–H2O gas mixtures [9, 10]. It was proved that the iron at the cementite/graphite interface is a reaction product from cementite decomposition [9].
In this report, the process of iron layer formation already reported in [9, 10] is further studied with emphasis on the effect of gas composition and temperature. The experiments were performed at both 600 and 700 °C in H2–CO–H2O gas mixtures. The results provide new information for the understanding of the phenomenon of iron layer formation during cementite decomposition and of the mechanism of metal dusting as well.
John Campbell , in Complete Casting Handbook (Second Edition), 2015
15.2.3 Olivine Sand
Olivine sand is sometimes used as an alternative to silica, primarily by foundries that cast high-alloy steels. Olivine possesses several advantages compared with silica, the most important being a lower and more uniform thermal expansion and inertness toward certain difficult alloys, notably manganese steels. Because olivine does not react with chromite, unlike silica, it can be tolerated at far higher levels in reclaimed chromite sand, and vice versa. Furthermore olivine with low loss on ignition (LOI), having a low serpentine content, is not regarded as being hazardous to health. Olivine sand does not react with alkaline binders. Some foundries use a 'dead-burnt' serpentine/olivine which is brown in colour and behaves similarly to ordinary olivine.
In spite of these valuable advantages, olivine is not widely used for two main reasons. The most important is that it is incompatible with the common binder system, furan no-bake (FNB), and requires specially formulated phenol-urethane no-bake (PUNB) binders. This means that olivine is used almost solely with alkaline binders and this leads to low sand reclamation rates with conventional equipment. A secondary disadvantage suffered by olivine sand is that it is, like chromite, made from crushed rock and thus possesses an angular grain shape that causes dusting during conveying and mixing. Finally, the refractoriness of olivine is related to its iron content which for steel casting should not be greater than 8% and for certain goods e.g. large carbon steel castings, maximum 6%.
Corrosion issues in water-cooled water-moderated energetic reactor (WWER) systems
K. Varga , ... J. Schunk , in Nuclear Corrosion Science and Engineering, 2012
Secondary systems (materials)
WWER secondary circuits were originally made mainly from carbon steel (with the exception of the stainless steel SG tubing and collectors), high alloy steel (turbine blades) or Admiralty brass of cupronickel alloy (MSR tubing, low pressure heater tubing and condenser tubing, although Loviisa, Temelin and Tianwan (Denisov et al., 1991) have titanium condenser tubing). Due to the extensive flow assisted corrosion (FAC) damage that occurred, which particularly affected the carbon steel high pressure heater tubing, moisture separator reheater tubing, wet steam lines and feedwater heater shells, the damaged sections were replaced with stainless steel equivalents. In addition, where possible, cupronickel components were replaced by copper-free ones so that the feedwater pH could be raised to suppress residual FAC damage and iron transport into the steam generators.
Stainless steel 08Cr18Ni10Ti used in WWER SGs has a slightly lower corrosion resistance, compared to Alloy 690 or even Alloy 800 but has much better corrosion behaviour under normal operating conditions than Alloy 600 and does not require as restrictive chemistry guidelines. In contrast to Alloy 600, which is very sensitive to stress corrosion in alkaline environments, stainless steel 08Cr18Ni10Ti is more sensitive to acidic conditions while a slightly alkaline environment is not detrimental and should be preferred if necessary.
At some stations it was found that the copper-based alloy condenser tubing suffered from numerous leaks, which increased impurity levels in the feedwater in spite of the presence of a full flow condensate polisher. To limit ingress and to permit operation at high pH the tubing at some East European WWER-440 plants has been replaced by titanium (Dukovany) or stainless steel (Paks and Kozloduy), allowing operation at high pH (Garbett, 2002). Other plants (Bohunice, Zaporozhe) that still operate with brass condensers have successfully tested and used amine-based AVT chemistry regimes (like ETA or morpholine) to minimize FAC problems.
At Temelin these changes were made before the station was commissioned and both units were completed with titanium condenser tubing and an all-ferrous secondary circuit. Titanium tubing and an all-ferrous circuit were also specified for the Tianwan units and similar secondary circuit materials will be specified for other new units, with the exception that the condenser tubes may be either titanium or stainless steel; no copper-based alloys will be used. The decision to use titanium tubing is expected to increase plant reliability and performance, as cooling water flow rates can be increased, which will not only improve condenser vacuum but will reduce deposition on to the condenser tubing (fouling), potential condenser leaks and permit the increase in secondary circuit pH to limit FAC damage and iron transport.
In Practical Machinery Management for Process Plants, 2005
Straightening Carbon Steel Shafts
Repair Techniques for Carbon Steel Shafts
For medium carbon steel shafts (0.30 to 0.50 carbon), three general methods of straightening the shaft are available. Shafts made of high alloy or stainless steel should not be straightened except on special instructions that can only be given for individual cases.
The Peening Method.
This consists of peening the concave side of the bend, lightly hitting it at the bend. This method is generally most satisfactory where shafts of small diameters are concerned—say shaft diameters of 4 in. (100 mm) or less. It is also the preferred—in many cases, the only—method of straightening shafts that are bent at the point where the shaft section is abruptly changed at fillets, ends of keyways, etc. By using a round end tool ground to about the same radius as the fillet and a 2 1/2-lb machinist's hammer, shafts that are bent in fillets can be straightened with hardly any marking on the shaft. Peening results in cold working of the metal, elongating the fibers surrounding the spot peened and setting up compression stresses that balance stresses in the opposite side of the shaft, thereby straightening the shaft. The peening method is the preferred method of straightening shafts bent by heavy shrink stresses that sometimes occur when shrinking turbine wheels on the shaft. Peening the shaft with a light (1/2 lb) peening hammer near the wheel will often stress-relieve the shrink stresses causing the bend without setting up balance stresses.
The Heating Method.
This consists of applying heat to the convex side of the bend. This method is generally the most satisfactory with large-diameter shafts—say 4 1/2 in. (∼ 112.5 mm) or more. It is also the preferred method of straightening shafts where the bend occurs in a constant diameter portion of the shaft—say between wheels. This is generally not applicable for shafts of small diameter or if the bend occurs at a region of rapidly changing shaft section. Because this method partially utilizes the compressive stresses set up by the weight of the rotor, its application is limited and care must be taken to properly support the shaft.
The shaft bend should be mapped and the shaft placed horizontally with the convex side of the bend placed on top. The shaft should be supported so that the convex side of the bend will have the maximum possible compression stress available from the weight of the rotor. For this reason, shafts having bends beyond the journals should be supported in lathe centers. Shafts with bends between the journals can usually be supported in the journals; however, if the end is close to the journal, it is preferable to support the shaft in centers so as to get the maximum possible compression stress at the convex side of the bend. In no event should the shaft be supported horizontally with the high spot on top and the support directly under the bend, since this will put tension stresses at the point to be heated, and heating will generally permanently increase the bend. Shafts can be straightened by not utilizing the compressive stress due to the weight of the rotor, but this method will be described later.
To straighten carbon steel shafts using the heating method, the shaft should be placed as just outlined and indicators placed on each side of the point to be heated. Heat should be quickly applied to a spot about two to three in. (∼50-75 mm) in diameter, using a welding tip of an oxyacetylene torch. Heat should be applied evenly and steadily. The indicators should be carefully watched until the bend in the shaft has about tripled its previous value. This may only require perhaps 3 to 30 seconds, so it really is very important to observe the indicators. The shaft should then be evenly cooled and indicated. If the bend has been reduced, repeat the procedure until the shaft has been straightened. If, however, no progress has been made, increase the heat bend as determined by the indicators in steps of about 0.010-0.020 in. (0.25–0.50 mm) or until the heated spot approaches a cherry red. If, using heat, results are not obtained on the third or fourth try, a different method must be tried.
The action of heat applied to straighten shafts is that the fibers surrounding the heated spot are placed in compression by the weight of the rotor, the compression due to expansion of the material diagonally opposite, and the resistance of the other fibers in the shaft. As the metal is heated, its compressive strength decreases so that ultimately the metal in the heated spot is given a permanent compression set. This makes the fibers on this side shorter and by tension they counterbalance tension stresses on the opposite side of the shaft, thereby straightening it.
The Heating and Cooling Method.
This method is especially applicable to large shafts that cannot be supported so as to get appreciable compressive stresses at the point of the bend. It consists of applying extreme cold—using dry ice—on the convex side of the bend and then quickly heating the concave side of the bend. This method is best used for straightening shaft ends beyond the journals or for large vertical shafts that are bent anywhere.
Here, the shaft side having the long fibers is artificially contracted by the application of cold. Then this sets up a tensile stress in the fibers on the opposite side which, when heated, lose their strength and are elongated at the point heated. This now sets up compressive stresses in the concave side that balance the compressive stresses in the opposite side. Indicators should also be used for this method of shaft straightening—first bending the shaft in the opposite direction from the initial bend, about twice the amount of the initial bend—by using dry ice on the convex side—and then quickly applying heat with an oxyacetylene torch to a small spot on the concave side.
Shafts of turbines and turbine-generator units have been successfully straightened by various methods. These include several 5,000-kw turbine-generator units, one 6,000-kw unit, and many smaller units. Manufacturers of turbines and other equipment have long used these straightening procedures, which have also been used by the U.S. Navy and others. With sufficient care, a shaft may be straightened to 0.0005 in. or less (0.001 in. or 0.025 mm total indicator reading). This is generally satisfactory.
0 Response to "What Is High Alloy Steel Made of"
Post a Comment