Microalloying of Steel Mikroiegiranje jekla F. Vodopivec, Inštitut za kovinske materiale in tehnologije, Lepi pot 11, 61001 Ljubljana Mechanisms of the effect of microalloying on strength and toughness of steels. Influence on grain size, precipitation hardening, processes of hot deformation, and economv of microallovina with Al Nb, i/, and Ti. Mehanizmi vpliva mikrolegiranja na trdnostne lastnosti in žilavost jekla. Vpliv na velikost zrn, izločilno utrditev, procesi vroče deformacije in gospodarnost mikrolegiranja z Al, Nb, V in Ti. 1 Introduction The microalloying of steel is a technology which has been intensively developed for about 25 years, and it exploits the theoretical know!edge on mechanisms of precipitation hardening, grain size control, and deformation of steel. The term "microalloying" is used because steels are alloyed with up to 0.05% of various elements, and an important influence on the following characteristics and properties is achieved: • austenite and ferrite grain size are diminished, and because of it yield stress, strength, and toughness are in-creased while the ductile/brittle fracture transition temperature is diminished; • precipitation hardening is achieved, this increases the yield stress and strength of steel, diminishes the toughness and increases the ductile/brittle fracture transition temperature; • the hardenability is improved and austenite/ferrite transition temperature is lowered; • the susceptibility of steel to strain ageing is eliminated; • the content of dissolved oxygen and sulphur in steel are diminished and the purity of steel is improved; • the shape and composition of non-metallic inclusions are changed and the isotropy of properties and the machinability of steel are improved; • the texture in non oriented electrical sheets is improved and the energy losses diminished. Microalloying elements are: aluminium, base element for steel deoxidation and for the decrease of oxygen is solution, niobium, titanium, vanadium, zirconium, boron, calcium, tellurium, antimony, tin, nitrogen, and in some cases also sulphur and lead. In this paper only microal-loying elements in the narrower sense will be discussed, i.e. those which influence the microstructure, strength and toughness of steel: aluminium, niobium, titanium, vanadium, and nitrogen which are in various combinations the basic constituents of high-strength structural steels and modem machine-building steels. In order to understand better the influence of microalloying elements on the mechanical properties and the hot vvorking process it is necessary to know the processes and reactions in steel involving these elements, and their compounds with nitrogen and carbon which form precipitates called in the following as disper-soide phases. The influenced processes are austenite grain growth, precipitation hardening in ferrite, and recrystalliza-tion of austenite during hot rolling. 2 Size and stability of austenite grains The first condition for the formation of small ferrite grains during the cooling of steel are small austenite grains and are obtained either by recrystallization of austenite after hot rolling at a relatively low temperature if a suitable delay of austenite grain grovvth is achieved during the hot vvorking, or during the cooling after normalization. Austenite grains grow through migration of boundaries, which can be hindered or stopped if the boundary is pinned to precipitates of dispersoide phases. When migration progresses, at first a concavity is formed at the precipitate, then the grain boundary envelopes it and finally bypasses it. This process requires an additional energy. The driving force for the growth is the tendency of material to reach a minimal total energy (E,) through the change of the shape and the size of grains, and is obtained by the minimal specific surface energy of grains. The total energy consists of the volume (E v) and the surface component (Ep). Ev is proportional to the grain volume, thus to D3, if D is linear dimension of grain, vvhile the surface component is proportional to D2. Schematically it can be written Es = I< D2 + I\\D3. The specific energy is thus: El = 1 D2 D + 1 E.g.: for D = 1, EJD = 2; for D = 2, E,/D = 1.5; for D = 3, Es/D = 1.33, etc. Thus total energy is the lower the coarses is the grain size. The prevention of the migration of a grain boundary is achieved when the distance among the precipitates is below a critical value. Instead of the distance betvveen the precipitates, which is difficult to be measured, the more easily measurable precipitation size (d) and volume part of dispersoide phase (/) are used in the analytical treatment of grain grovvth. The relationship between the grain size— D, the volume part of precipitates—/, and their size—d is according to Zener1 given by the equation: D _ 4d J ~ 3f The above equation vvas further developed for the grovvth of austenite grains in structural steel by Gladmann and Pickering2. They have assumed that in grain grovvth the energy D \Z 2 J is released. In the equation S—represents the boundary migration, y—the surface energy of austenite, and Z—the ratio betvveen the size of a growing grain and the average grain size in the matrix. It is evident that the grain growth is possible only if Z > 4/3, otherwise the growth energy change is positive and a spontaneous process is not possible. An exception represents the cases of very great growth driving force, e.g. a grain shapes which greatly deviates from the equilibrium. The equation was further transformed into the expression eonnecting the critical size of precipitates, c/j., with other easily nteasurable parameters: d i-. The equation shows that grains grow with the growth of the critical size and the dccrease of the content of precipitates, as vvell as with the increasing non-uniformity of grain size. A spontaneous grovvth is initiated the easier the greater is the initial non-uniformity of austenite grain size. For better understanding it can be mentioned that after one-hour of heating of a Cr-Ni carburizing steel at 920°C, i.e. before anormal grain grovvth, the ratio of maximal and minimal grain size is Z = 3.18. By the same quantity of the dispersoide phase the precipitates are the more efficient the smaller is their size, i.e. the greater is their volume density and thus the smaller is their mutual distance. Precipitates are not completely stable at the grain grovvth temperature and grow at prolonged an-nealing tinte and especially at higher temperatures loosing the pinning efficiency. In structural steel the precipitates of sizes below 10 nni represent a low hindrance for grain grovvth because of their instability caused by the high ratio of surface to the total energy. Efficient precipitates are formed by dispersoide phases vvhich are dissolved in austenite at the heating of steel before the hot rolling or forging. As steels are becoming cool, the solubility of dispersoide phase is diminished, and precipitates are formed vvith size depending on the temperature and the length of isothcrmal annealing. During the transformation and the recrystallization the grovvth rate of ali grains is not uniform, single grains grovv faster, reach a lovver total energy, become more stable, and at sufficient temperature grovv at the expense of their neigh-bours. This is the explanation why a microstructure of grains of different size is found in normalized steel vvith a too lovv quantity of the precipitates for complete pinning of the migration of austenite grain boundaries. The grains can grovv also by coalescence if their space orientation is similar and are parted by lovv-anglc boundaries. This occurs in textured materials. The boundary migrating at the extent of a neighbour grain is concave. A simplificd explanation is that the atoms on the concave side are on average more tinte bound in the crystal lattice. The migration of crystal boundary is produced by the difference in the number of atoms vvhich are deplaced over the grain boundary because of thermal oscillation. The number of jumps front the convex to the concave side of the boundary is equal to the number of jumps in the reverse dircction, but on the concave side more of oscillating atoms are retained. This produces a flovv of atoms from the convex to the concave side, i.e. the shift of crystal boundary in the opposite dircction. The theoretical cxplanation for the migration process of a crystal boundary tovvards the centre of curvature is found in ref.3, vvhere it is suggested that the driving force for the boundary migration is the decrease of surface energy. The rate of migration is described by a parabola of the form D = Do + A'1/n vvith D0—an initial grain size, D—the grain size after an annealing time t, and n—the grovvth exponent. Theoreti-cally n is 2 vvhile empirically the values betvveen 2 and 4 vvere measured. The connection betvveen the grain size (D) and the yield stress (Re) is given by the Hall-Petch equation Re = Re0 + KD-1'2 vvith Re0 as a constant depending on the contposition and the microstructure of steel . The constant K is a measure for the hindering effect of crystal boundaries on the mobility of dislocations. In Fig. 1 taken from the ref.5 the relation betvveen the grain size, expressed by D1?2 and the ASTM number, and the yield stress of steel vvith 0.17% C and 0.8% Mn is shovvn. The decrease of the grain size from ASTM number 5 to ASTM number 10, obtained through the microalloying produces an increases in yie!d stress of steel for about 50%. This increase takes plače at an increased toughness and a decreased ductile/brittle fracture transition temperature (Tp) at notch toughness test. The proposed relation is ^r = T0 + K D -1/2 with To and K constants depending on the contposition and the microstructure of the steel4. Figure 1. Relation betvveen the grain size expressed as D , the ASTM number, and the yield stress of steel vvith 0.17% C and O.S^i. Mn (Ref.5). Slika 1. Odvisnost med velikostjo zm izraženo kot D-1/2 in ASTM razredom ter mejo plastičnosti jekla z 0.17% C in 0.8% Mn. Po viru5. The movement of dislocations in the lattice is hindered by the Peierls-Nabarro force (rpn), and each crystal bound-ary produces an additional obstacle for the movement. The 5 10 15 Grain size d~!/J v mm 10 11 12 13 ASTM number kp/mm mm"1'' total force (Ts) essential for the movement of dislocations in polycrystalline material is: Ts = Tpn + K D~1'2 A piling-up of dislocations at the grain boundary is re-quired in order to accumulate a sufficient driving force for the penetration of dislocation into the neighbour grain with a different space orientation6. 3 Dispersoide phases Dispersoide phases are carhides and nitrides, very fre-quently also carbonitrides since carbides and nitrides of mi-croalloying elements are mutually completely soluble. The composition of dispersoides depcnds on the amounts of mi-croalloying elements, nitrogen, and carbon in steel. If the content of microalloying elements, nitrogen or carbon is too high, dispersoide phases can form already in the melt or during the solidification of steel. The size of precipitates in this čase is 100 nm or more, accordingly small is their volume density, and low the hindering effect at standard size of austenite grains. In microalloyed steel in which the content of microalloying elcmenLs for the most part does not exceed 0.05%, the sufficient volume density of precipitates is not achieved if they are formed in the melt or during the solidification. In this čase they are enriched on the solidification interfaces or in eutectic clusters, vvhich decreases the ductility of the steel. Such example represent A1N and Nb(CN) formed during the solidification of steel manufactured in electric are furnace7, vvith a high content of nitrogen, aluminium, and niobium. The mechanism of grovvth of precipitates involves the solution of small particles vvith a greater specific surface energy and the diffusion of microalloying components on coarse, more stable particles. The grovvth of precipitates, often named as Ostvvald ripening, is deseribed by the Wagner8 equation df ~ d30 = 16 aDCV 9 RT -t vvith dt precipitate diameter after the annealing time t do precipitate diameter in time 0 u surface tension betvveen the matrix and pre- cipitate D diffusivity of constitutive atoms C concentration of constitutive atoms V molar volume of precipitate R gas constant T absolute temperature The equation shovvs that the rate of precipitate grovvth will be at constant other conditions the faster, the faster is the diffusivity, the greater is the concentration of constitu-tive atoms in solution, and the higher is temperature. Thus the ideal dispersoide is that vvith the lovvest solubility of constituents, and vvith the lovvest diffusivity of microalloy-ing element, since the diffusivity of nitrogen and carbon in interstitial solution is very fast. Let us assume that the steel contains 0.03% A1N vvhich ensures the austenite grains size after normalization of 6-7 ASTM number9. Fig. 2 piesents the calculated A1N precipitate size for such a steel after one hour fiolding at various temperatures, the content of aluminium nitride, the volume density of precipitates, and the relative austenite grain size. The solubility product used for the calculation is in good agreement vvith the A1N solubility determined for Cr-Ni carburizing steel9. If the heating temperature of steel is inereased from 900 to 950° C, the same hindering effect can be obtained vvith an about 40% higher content of nitride, vvhile above 1000°C the pinning effect of aluminium nitride is very rapidly diminished. 0.15 i.0,10 300 200 100 ,0.05 OL o r 30 1000 1100 Temperature °C Figure 2. Relation betvveen the annealing temperature and the precipitate size. the number of precipitates per unity of volume, the content of A1N in solution, and the austenite grain size. The theoretical AIN content is 0.03%. The calculation is based on the A1N solubility product given in ref.2. Slika 2. Odvisnost med temperaturo žarjenja in velikostjo izločkov, številom izločkov na enoto prostornine, količino AIN v raztopini in velikostjo zm austenita. Teoretična vsebnost AIN 0.03%. Izračun je izvršen na osnovi topnostnega produkta za AIN v viru2. In microalloyed steel usually there are 2 or 3 grain grovvth inhibitors; AIN, Nb(CN), TiC and VN. The presence of precipitates formed by the addition of 0.03% niobium to the steel vvith 0.10-0.20% C ensures grain sizes of ASTM number 10-11 after normalization. The most frequent dispersoide is aluminium nitride (AIN) vvhich is found in ali steels deoxidized, and thus mi-croalloyed vvith aluminium. The solubility of AIN and of other dispersoide phases in austenite in struetural steels is given by the solubility product. Ref.10 gives the follovving solubility product for aluminium nitride log(Al x N) = -6770/T + 1.48 In the above equation N and Al represent the vveight content of both elements in solution in the steel, and T is the temperature in K. According to ref. 10, 11, and solubility products for other dispersoide phases are 12 the log(Ti x C) = — 10475/T + 5.33 ti log(Ti x N) = - 8000/T + 0.32 12 log (V x C) = - 9500/7" + 6.72 11 log(V x N) = - 8330/T + 3.46 10 log(Nb x C + N) = - 6770/7" -f 2.26 10 In some references also other equation for the solubility of dispersoides is found but they do not differ significantly from the above given. The solubility of ali the dispersoides, but of vanadium carbide, in austenite with up to 0.2% C and 0.01% N is small and at the normalizing temperature less than 10% of the quantity at the temperature of about 1200° C. On the contrary, the solubility of vanadium carbide in austenite is very high, and this dispersoide is completely dissolved already at about 900°C in steel with 0.2% C. Other dispersoides are very stable because of the low solubility at the normalizing temperatures and the inhibition of grain growth is diminished only at higher temperatures. During the cooling from the solubility temperature and at isothermal holding during such cooling the formation of precipitates is very slow (Fig. 3) due to slow formation of nuclea though the solid solution is highly oversaturated13. The explanation for the slow formation of nuclea is the great dilution since the content seldont exceeds 0.03% vvhich e.g. represents 3 atoms of titanium per 10000 atoms of iron. The number of atoms of microalloying elements is thus very lovv and consequently the rate of formation of sufficient statistic aggregations of atoms from vvhich precipitation nuclea are formed is very slovv. The kinetics of the precipitation during the holding after direct cooling from the solubility temperature is a slovv parabola highly different from that describing the formation of precipitates in austenite quenched from the solubility temperature and then reheated (Fig. 3). The kinetics of A1N formation is in this čase a step parabola indi-cating that the rate of grovvth of precipitates is determined by the diffusion rate of aluminium on the nuclea formed during the reheating of steel, due to the high oversatura-tion because of the cooling to ambient temperature or to the double transition of the austenite/ferrite phase boundary on vvhich the solubility of A1N is changed strongly. dium (Fig. 4). The specific vveight of niobium carbide is higher than that of aluminium nitride, the vveight solubility of both in austenite is similar, thus the same vveight content of niobium in austenite gives less precipitates. Conse-quently, if seems probable that niobium hinders the migra-tion of boundaries also by some other mechanism, e.g. by a segregations on grain boundaries vvhich produces a greater number of precipitates on these boundaries as it could be expected from the average niobium content in steel. Ref.14 presents micrographies shovving that the boundaries or sub-boundaries of austenite grains are marked vvith strings of precipitates vvhich confirm the possibility of an intercrys-talline segregation of niobium. The size of austenite grains is thus related to the thermal deformation history of steel. Ref.15 describes a bimodal size distribution of precipitate after the rolling of niobium steel from 1050°C vvhich can also be explained supposing an intercrystalline segregation of niobium. 30 o. u o 20 10 \ Č 4320 ~~--- \ \ ^^ -V i * - —a- Nb 0,05 0.10 Content of Nb and V . % 0.15 Figure 4. Relation betvveen the amounts of niobium and vanadium in steel and the size of austenite grain after half-hour and 8-hour austenitizing at 920°C. Basic steel composition: 0.18% C, 0.95% Mn, 0.28% Si, 1.0% Cr and, belovv 0.002% Al (Ref.16). Slika 4. Odvisnost med količino niobija in vanadija v jeklu in velikostjo zm austenita po polurni in 8 umi austenitizaciji pri 920°C. Osnovna sestava jekel: 0.18% C, 0.05% Mn, 0.28% Si, 1.0% Cr. pod 0.002% Al,. Po viru16. Figure 3. Kinetics of A1N precipitation after various thermal history of steel vvith 0.11% C, 0.49% Mn, 0.029% Al, and 0.0063% N. Slika 3. Kinetika precitipacije A1N po različni temiični zgodovini jekla z. 0.11% C. 0.49% Mn, 0.029% Al in 0.0063% N. It must be mentioned that niobium if its concentralion exceeds about 0.035% and at high nitrogen contents—the limit is at about 0.012%, is bound during the solidification process into a carbonitride very rich in nitrogen and practi-cally unsoluble during heating the steel before the rolling7. Niobium bound in this phase is lost as active microalloy-ing element, thus the microalloying vvith niobium in steel molten in electric are fumace is economical only up to about 0.03%. In CrMn čase hardening steel by already 0.02% Nb the same stability and size of austenite grains is achieved as vvith the same amount of aluminium or vvith 0.1% vana- 4 Microalloying and precipitation hardening The precipitation hardening is one form of dispersion hardening, i.e. hardening caused by a nevv phase vvhich is found in small quantities in the metallic matrix. The general ex-pression describing the relations betvveen the quantity of precipitates (/), their size (d), the shear modulus (G), the Burgers veetor of dislocations (b), and inerease of strength (ARt) was proposed by Hombogen17 in the follovving form ARr = A-921 d K is a constant vvith a value I\ = 1 for a polycrystalline material and uniformly distributed spheroidal particles of the nevv phase. The hardening is proportional to the third root of the quantity of precipitated phase and inversely proportional to the particle size of that phase. It is thus more strongly dependant on the size than on the quantity of precipitates. The precipitation hardening is stable only till the shear moduls is not diminished because of the temperature change or the precipitates do not hinder the moving of dislocations. In microalloyed steel the precipitates formed at 0.029AI, 0,0063 N 1300Nb), d =0.C )5pm with Nb as niobium content in vveight %, d—the size of niobium carbonitride precipitates, and ARe—the increase of yield stress. The exponent at the niobium content proves that the expression was derived through simplification of the Hornbogen equation. A disadvantage of the expression is the lack of parameters considering the temperature of formation of precipitates and the shear modulus, thus it can be used only for heat treatment by quenching and ageing at a selected temperature. Fig. 5 presents the influence of precipitates size at constant niobium content, and the content of niobium at constant size of 50 nm precipitates on the hardening effect. Already a small amount of niobium is efficient if present in steel in small precipitates. The increase of the content of niobium does not improve the hardening effect to an economica!ly justified extent. Fig. 5 further proves that precipitates of an average size of 25 nm, vvhich can be found in microalloyed steel after normalizing, cause hardly a hardening effect. Practically only a part of precipitation hardening effect can be industrially exploited hovvever it is not negligible 0 0.01 0.02 Precipitates size d. pm t_t_i_i 0 0,02 0,04 0,06 Content of Nb,% Figure 5. Relation betvveen the size of precipitates in steel vvith 0.03% Nb or NbC concentration in 5 nm precipitates, and the increase of yield stress. Slika 5. Odvisnost med velikostjo izločkov v jeklu z 0.03% Nb oz. količino NbC v izločkih z velikostjo 5 nm in povečanjem meje plastičnosti. from the vievvpoint of the material strength. E.g. a small change in basic composition of the steel vvith a yield stress above 350 N/mm2, and microalloying vvith niobium and vanadium can give a yield stress above 470 N/mm2, vvhere precipitation hardening due to formation of vanadium carbide during the cooling of steel after normalizing contributes for about 50 N/mm2 . As already mentioned, the precipitates formed in the approximate temperature range 570 to 620° C are efficient. At lovver temperatures the formation of precipitates is too slovv due to the slovv diffusion of vanadium, and it could be exploited only by a longer annealing or in a very slovv cooling vvhich is economic only in coils. The effects of the diminution of grain size and of the precipitation hardening on the yield stress are additive, their influence on the other tvvo very important properties, the notch toughness and the ductile/brittle fracture transition temperature, is opposite. Diminished grain size increases the toughness and decreases the transition temperature vvhile the precipitation hardening has an opposite effect. Fig. 6 presents, according to data in ref.4, some relations vvhich confirm the above conclusions for a standard as normalized Nb-V microalloyed steel. By thermal treatment, e.g. by normalizing and through the rate of cooling, a rather different relation betvveen the yield stress and the toughness transition temperature can be achieved, even an unaxcept-able transition temperature can be obtained vvhich nullifies ali the advantages of microalloying. 5 Microalloying and hot deformation Most steel products are manufactured by hot rolling vvhen the steel cross section is reduced from pass to pass at drop-ping temperature till a final thickness of plate, strip or bar is obtained. A similar sequence of events is found by forgirig only the sequence of partial deformations is less uniform. During the rolling process the steel is cooled partially by radiation and the convection into the surroundings, and par-tially by contact vvith the cooling vvater, rolls, hammers or 400 300 E E ■ 200 tr 100 100 ARe N/mm2 : 17,36 + 1.81 D"v2 Re=f(d) 250 r+80 VPT= f (IU) TT = f (D) 0,15 C, 0,46 Si. 1,16 Mn , 0 036 V, 0,015 Nb normahsed L_I________i 12 10 6 ASTM number 4 +60 +40 +20 -4 o X 20 -I -40 o o O) i_ ZD a O) CL E (D C O Cn c a -60 20 40 60 80 Gram size D, jjm 100 Relationships grain size - yield stress grain size - transition temperature precipitation hardening - - transition temperature Figure 6. Relation between the grain size in as normalized steel, the precipitation hardening, and the yield stress increase due to the precipitation hardening (ARe), and the notch toughness transition temperature (PT). Slika 6. Odvisnost med velikostjo kristalnih zm v normaliziranem jeklu, oz. izločilno utrditvijo in mejo plastičnosti povečano zaradi izločilne utrditve (ARe) ter prehodno temperaturo žilavosti (PT). other cooler parts of equipnient. Because of the successive deformations the steel is in deformed state for some time and it contains a great number of point and line defects. The rate of diffusion processes in the deformed ntatrix is very fast. Jonas and covvorkers21 22 found that the nucleation rate of precipitates was for an order of magnitude faster during the deformation, and that the rate of precipitates growth in deformed austenite was for two orders and during the deformation even for three orders of magnitude greater than in not deformed or in recrystallized austenite. The static recrystallization vvhich eliminates from austenite the deformation energy delayed for a few seconds corresponds thus to a 100 or even 1000 sec. long annealing. Any component which delays the recrystallization thus highly accelerates the proeess of precipitaton but only as long as austenite remains unrecrystallized. As soon as the recrystallization is finished the rate of precipitation is diminished again. E.g. in labo-ratory rolling of 12 mm plates from 60 mm billets the steel remained between the rolls for 0.47 seconds, and total time of rolling was 70 seconds. Let us assume a great rate and an uniform precipitation in the period when steel is deformed betvveen the rolls. The rate of precipitates growths is de-seribed by approximate cubic parabola which can be simpli-fied for a rough evaluation into the expression t/'/'5 % Kt, t being the time. The calculation shows that the ratio of precipitates size in unrecrystallized austenite (dan) and in recrystallized austenite (dar) is dan/dar = 4.5. If steel is rolled by uncompleted intetpass recrystallization and if it contains a small quantity of microalloying elements the unhomogenity of microstructure represented by the number of anormally grown grains is the greater the lower is the rolling temperature9,2 because of the unhomogenity of precipitation during the rolling. During the rolling of low and medium alloyed steel with an austenite microstructure static recrystallizaton is the basic proeess for the elimina-tion of deformation energy. Dynamic softening processes and static recovery are virtually negligible. In absence of interpass recrystallization static recovery rapidly eliminates the deformation hardening, and it does not change the size of austenite grains. Niobium is the microalloying element which has the strongest delaying effect on the rate of static recrystallization of austenite. Two explanations are pro-posed for the mechanism of the effect of niobium. Tite first elaims that the effect is linked to niobium in solid solution in austenite24. The proeess of static recrystalization is initiated on the grain boundaries, it seems thus that the inhi-bition of formation of recrystallization nuelea on boundaries is connected to the presence of niobium at these boundaries. Fig. 7 shows that the temperature of completed interpass static recrystallization of austenite is already by 0.02% niobium inereased for about 100° C in the CrMn carburizing steel. The weight content of 0.02% of niobium means that the solution tn austenite contains appr. 1.1 niobium atom per 104 iron atoms. A logic conclusion is that sueh a dilution could hardly influence the proeess linked to shifts of iron atoms and it seems justified to conclude that the austenite grains boundaries are richer in niobium due to a segregation. This explanation is supported also by the fact that small amounts of niobium improve the nardenability of steel through the delaying the nucleation of ferrite below the transformation temperature. Thus niobium can hinder the formation of recrystallization nuelea and ferrite by a similar mechanism. o: CrMn cc ise harder R o r ing steel u - O O 6 G- / i P —m— / / / /' /F / ^__ 6 E E in cz 4 o t_ OT O!-o*a-U---. V - ---------10 1200 1100 1000 900 800 Initial rolling temperature, °C Figure 7. Influence of the initial rolling temperature on the ratio length/width (R) and on the number of unrecrystallized austenite grains (P). Steel G: 0.14% C. 1% Mn, 0.85% Cr, 0.02% Nb, and 0.0078% N; steel F: 0.16% C, 1.1% Mn, 0.98%: Cr, 0.025% Al, and 0.0095% N. Slika 7. Vpliv začetne temperature valjanja na razmerje dolžina/širina (R) in na število nerekristaliziranih zm austenita (P). Jeklo G: 0.14% C, 1% Mn, 0.85% Cr, 0.02% Nb in 0.0078% N; jeklo F: 0.16% C, 1.1% Mn, 0.98% Cr, 0.025% Al tn 0.0095% N. The second hypothesis links the influence of niobium on the recrystallization on precipitates formed during the deformation. Two questions are not explained by this hy- pothesis: why the precipitates of other microalloying elements, e.g. TiC, and VN and A1N, vvhich are also formed during the deformation, hinder the static recrystallization of austenite to a much lesser extent (Fig. 8), and why the recrystallization process takes plače when the content of niobium in solid solution is diminished below a limit of about 0.005% due to the formation of precipitates. Both explanations of the effect of niobium are found in recent papers on microalloying, and it is left to the reader to chose the more probable significance weighting the significance of empirical findings. 300 0,050 0,100 0.150 0.200 Content of elements.0/« 0.250 Figure 8. Influence of the content of various microalloying elements on the hindering temperature of static recrystallization of austenite. Slika 8. Vpliv vsebnosti raznih mikrolegimih elementov na temperaturo zaustavitve statične rekristalizacije austenita. In Fig. 9 the effect of rolling temperature on the content of A1N and NbC in various steels, and on the ferrite grain size given as intercept length by rolling 15 mm plates from 60 mm billets in 7 passes is shown. Ali the steels were heated to 1200°C before the rolling. In steel vvith-out niobium where the interpass recrystallization of austenite is fast and complete, only few precipitates are formed during the rolling and the influence of temperature on the arnount of precipitates is hardly perceivable because at de-creasing rolling temperature the content of A1N formed during the rolling is very slowly inereased. The precipitation behaviour in niobium steel is significantly different because austenite remains between passes for longer time unrecrys-tallized, or the quantity of austenite which does not re-crystallize at ali betvveen passes is inereased, and thus the precipitation is faster. Below a limit temperature austenite remains completely unrecrystallized between passes and the precipitation is accelerated to sueh extent that practically ali A1N and NbC are precipitated in the relatively short rolling time of 1 minute. On the base of the processes of recrystallization and of precipitation tvvo technologies of rolling of microalloyed steel were developed. In thermomečhanical rolling the slabs are rolled to a thickness which is 30-50% greater than the final thickness of plates, the rolling is stopped till steel temperature drops belovv about 950°C, and then the plates are rolled to the final thickness in scveral passes, the number de-pends on the strength of the rolling stand, and cooled in air. Defomied austenite is during the eooling very rapidly trans-formed into finegrained ferrite and pearlite25, while A IN and NbC precipitates hinder the growth of ferrite grains after the transformation. A finegrained microstructure with high strength and toughness is obtained, and i! supports the precipitation hardening with VC during the air eooling of plates with an acceptable deteriorating effect on noteh toughness. 0 _L _L 1200 1100 Initial rolling 1000 900 temperature , °C Figure 9. Relation betvveen the initial temperature of tvvo steel vvith a similar basic composition, one microalloyed vvith niobium, the amounts of A1N and NbC precipitated during the rolling and the grain size after air eooling. Slika 9. Odvisnost med začetno temperaturo dveh jekel s podobno osnovno sestavo, eno pa mikrolegirano z niobijem na količino A1N in NbC, ki sta nastala med valjanjem in na velikost zrn po ohladitvi na zraku. This technology is applied for steel microalloyed vvith niobium and vanadium vvhich can achieve yield stresses up to 500 N/mm2 at earbon contents belovv 0.15%. Similar properties are obtained vvith the combination of rolling in less controlled temperature range and normalization anneal-ing. Fig. 10 represents the various hardening mechanisms in normalized steel microalloyed vvith aluminium, niobium, and vanadium. The advantages of microalloying can be exploited also through the rolling process vvith controlled recrystallization. This method demands an exact harmonization of steel composition vvith the degree of deformation and the pass se-quence, since the temperature and the per pass deformation must enable the complete interpass recrystallization, and si-multaneously also the formation of precipitates vvhich hinder the grovvth of recrystallized austenite grains. Fig. 11 presents mechanical properties of three steel of similar composition vvhich vvere rolled under conditions of controlled recrystallization. In both microalloyed steels much better properties are achieved than in the comparative steel dovvn to about 800° C when the transformation of austenite during the rolling, the deformation hardening, and the formation of texture during rolling start to occur. By the same rolling conditions the microstructure of vanadium steel is more ho-mogeneous because of less of microstructural unhomogene-ity due to the or incompleted interpass recrystallization. At stili lovver rolling temperatures the deformation hardening, 20mm plate from steel with 0,18 C ; 0.4 Si , UMn ; 0.02P , 0.025AI, 0.0042Nb 0,06V, 0.12Cr , 0,21 Cu , 0,10Ni, normalised yield stress 488N/mm2, grain size ASTM numberl! 500 VC NbC Al N Cu , Cr, Ni. P. Si precipitation hardening diminution of grain size substitution hardening 10 7. 9 K 9 5 24 C -perlite " C and N interstitlaf hardening n Natural yield stress and y uncontrolled irnpurities Figure 10. Constitution of yield stress in a 20 mm sheet of normalized microalloyed steel with 0.18% C, 0.4% Si, 1.4% Mn, 0.025% Al, 0.042% Nb, 0.06% V, 0.12% Cr, 0.21% Cu, and 0.10% Ni. Slika 10. Zgradba meje plastičnosti v 20 mm pločevini iz normaliziranega mikrolegiranega jekla z 0.18% C, 0.4% Si, 1.4% Mn. 0.025% Al, 0.042% Nb, 0.06% V. 0.12% Cr, 0.21% Cu in 0.10%- Ni. microstructural nonhomogeneities, and strain anisotropy in-crease, therefore a too low finishing temperature has a harm-ful effect. An even higher strength can be achieved by a proper coiling temperature since slow cooling enables a greater precipitation hardening. 6 Economy of microalloying Microalloying is the more efficient the more dispersoide is dissolved at heating before hot vvorking. The quantity of dissolved dispersoide is the greater the closer are the concentrations of the constituting elements to the stoichio-metric ratio; e.g. the amount of dissolved A1N in austenite will be the greatest if the weight contents of aluminium and nitrogen in steel are 2 : 1. A too great deviation of one or the other element causes that less dispersoide is dissolved in austenite, and less precipitates are formed during the rolling and the cooling. This explains why at high aluminium contents, over 0.04%, austenite grains are coarser and less stable than at a lower content of aluminium and at the same content of nitrogen about 0.01 %. For the stability of austenite grains the nature of precipitates is not important, only their number and stability matter, or more correctly, the number of precipitates per unity of volume of austenite. Theoretically the presence of about 0.03% A1N or of a corresponding quantity of other dispersoides of precipitates of equal size is needed to at-tain a sufficient stability of austenite grains. The contents of microalloying elements and of dispersoide phases giving equal volume densities of precipitates are given in Table 1 for the most frequent dispersoides. Aluminium assures a sufficient density of precipitates already at the lovvest content, while the highest content is required in the čase of vanadium carbide since this carbide is much more soluble E E cn c o; i_ i/i a> 10 c