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Dual phase steels

Dual phase steels






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    Dual phase steels Dual phase steels Document Transcript

    • ANNUAL REVIEWS Further Ann. Rev. Mater. Sci. 1981. 11:245-66 Quick links to online content Copyright © 1981 by Annual Reviews Inc. All rights reserved DUAL PHASE STEELS +8665 M. S. Rashidby University of Tennessee - Knoxville - Hodges Library on 05/22/12. For personal use only. Metallurgy Department, General Motors Research Laboratories, Warren, Annu. Rev. Mater. Sci. 1981.11:245-266. Downloaded from www.annualreviews.org Michigan 48090 INTRODUCTION Dual phase steels are characterized by a microstructure conslstmg of a bout 75-85 vol% ferrite (a-iron) with the remainder being a mixture of martensite, lower bainite, and retained austenite (Figure 1). The name "dual phase" was coined in the mid 1970s to describe ferrite-martensite microstructures (I), but dual phase steels usually contain more than the two phases implied by their name. They are essentially low carbon steels that are thermomechanically processed to have better formability than ferrite-pearlite steels of similar tensile strength. The stress-strain behavior of dual phase steels (2) is characteristically different from that of ferrite-pearlite steels such as plain carbon steel or the microalloyed, high strength low alloy (HSLA) steels (Figure 2). The ferrite-pearlite steels have yield point elongation, a high ratio of yield strength to ultimate tensile strength (YSIUTS), and their strength and ductility (uniform elongation) are inversely related (Figure 3). Dual phase steels have a continuous stress-strain curve with no yield point elonga­ tion. They work-harden very rapidly at low strains, have a low YS, a high UTS and hence a low YS/UTS ratio. They have better formability than the ferrite-pearlite steels of equivalent tensile strength and their strength­ ductility data fall on a separate curve (Figure 3) than that for ferrite­ pearlite steels. Ferrite-martensite steels were developed in the mid 1960s concurrently at BISRA (British Iron and Steel Research Association) in the United Kingdom (3-5) and Inland Steel Corporation in the US (6). Both efforts were concerned mainly with producing steels for tinplate by drastically quenching low carbon steels from temperatures close to the critical temperature. The BISRA objective was to develop a compact annealing process for making steels of about 500 MPa tensile strength; Inland Steel effort was directed toward producing steels with about 1000 MPa tensile 245 0084-6600/81/080 1-0245$0 1.00
    • 246 RASHID Transformation Productby University of Tennessee - Knoxville - Hodges Library on 05/22/12. For personal use only. Annu. Rev. Mater. Sci. 1981.11:245-266. Downloaded from www.annualreviews.org Ferrit 10 m Figure J Scanning electron micrograph of a dual phase steel. The microstructure consists of a fine-grained ferrite matrix with a uniform distribution of about 20 volume percent transformation product , which consists of martensite, retained austenite and bainite (2). 980X 600 GM ----- -----� , Plain Carbon I-+---Totul Elongotion, et ----I i o I 10 -11-0.2 Percent Stroin Figure 2 Schematic stress-strain curves for plain carbon, HSLA, and dual phase steels. SAE 950X and 980X are Society of Automotive Engineers designations for HSLA steels of different s trength levels. GM 980X is a General Motors developed dual phase steel. GM 980X is more ductile than SAE 980 X although both steels have similar tensile strength.
    • DUAL PHASE STEELS 247 700 3OX o ... � 600 of III C � f 1 R Ferrite-pearliteby University of Tennessee - Knoxville - Hodges Library on 05/22/12. For personal use only. .!! 500 . ;; c Annu. Rev. Mater. Sci. 1981.11:245-266. Downloaded from www.annualreviews.org ,! SAE 950 .!! o � 400 , 5 " " Figure 3 Strength-ductility relationship Plain Carbon � of dual phase steels compared with that for plain carbon and HSLA steels. The dual phase steel curve is far above that Uniform Elongation, 70 for ferrite-pearlite steels (40). strength. The BISRA process involved intercritical heating followed by quenching either into water with post tempering or into low temperature salt or liquid metal baths, the bath serving as the tempering medium. Although the first Inland Steel experiments were directed toward a fully quenched martensitic product, subsequent work was also conducted on intercritically heated and quenched sheet. Neither approach recognized or researched the potential for improved formability in these steels. Grange (7, 8) later investigated methods for producing fibered ferrite­ martensite microstructures. This process requires cold rolling a pearlitic structure to produce elongated patches of pearlite, which after a subse­ quent intercritical heating and quenching treatment produces fibers of martensite in a ferrite matrix. Steel having such microstructures has a better combination of strength and toughness than the same steel before the thermomechanical treatment. Although indications of better forma­ bility were evident in these data, the importance of this aspect was not elaborated. Bailey researched ferrite-martensite steels (9, 10) in the mid 1970s and reported that the strength of low carbon steels can be increased at the expense of ductility by intercritical heating followed by quenching and tempering. The ductility of these "dual phase" steels was comparable to that of ferrite-pearlite steels of similar tensile strength. Like some of the
    • 248 RASHID earlier work, this study also emphasized strength rather than formability improvement but it dealt with steels thicker (up to 2 mm) than tinplate. The development of dual phase steels that were tailored for improved formability was triggered in the early 1970s by conflicting demands in the US automotive industry for reduced weight to increase fuel economy, and increased weight to satisfy newly imposed safety and ecological standards. Steels were sought with strength-ductility combinations sub­ stantially better than existing grades of high strength sheet steels toby University of Tennessee - Knoxville - Hodges Library on 05/22/12. For personal use only. enable the fabrication of complexly shaped automotive components using Annu. Rev. Mater. Sci. 1981.11:245-266. Downloaded from www.annualreviews.org techniques perfected for plain carbon steel. Concurrent development of the desired steels occurred in Japan and the US. Matsuoka & Yamamori ( 1 1) and Hayami & Furukawa ( 1) reported the development of intercritically annealed, microalloy-free (vanadium, niobium and titanium) Si-Mn steels, while Rashid (2, 12, 13) reported on the development of intercritically annealed, microalloyed dual phase steels. These steels had far better formability than any of the previously reported high strength sheet steels and represented a breakthrough in high strength steel development. Bucher & Hamburg demonstrated that the research data reported by Rashid could be dupli­ cated in commercially produced steels ( 14, 15). Later, Coldren & Tither ( 16) showed that dual phase steels can also be produced as-rolled directly off the hot mill. Subsequently, numerous researchers have reported on variations of these approaches using modified steel compositions ( l7-20). The various developments made it clear that a dual phase microstructure by itself did not automatically guarantee good formability, but that the objective of good formability combined with high strength can be accom­ plished by proper control of steel composition and process variables. Low to high carbon steels with martensite-austenite microstructure, with martensite being the major phase, have sometimes also been referred to as "dual phase" steels (2 1 22) although they do not satisfy the , strength-ductility criteria shown in Figure 3. Martensite-austenite steels have unique combinations of mechanical properties but do not fall into the category of "dual phase" steels described earlier and should be dealt with separately. Publications and research on the highly formable dual phase steels have increased exponentially ( 11-20) since their inception just a few years ago, and a unified treatise of published literature is attempted in this article. This review summarizes the historical development of the steels and describes the present state of the art of the various approaches that have been used to produce the steels. Current understanding of the transformation and deformation mechanisms are discussed and the strength ductility relationships in these new steels are briefly reviewed.
    • DUAL PHASE STEELS 249 PROCESSING METHODS Dual phase steels have been produced by continuous annealing, as-rolled directly off the hot mill, and by batch annealing. Considerably more research and production activity is reported in continuous annealing than in the latter two approaches. Continuous Annealedby University of Tennessee - Knoxville - Hodges Library on 05/22/12. For personal use only. All reported continuous annealing processes have three common salient Annu. Rev. Mater. Sci. 1981.11:245-266. Downloaded from www.annualreviews.org features (Figure 4), namely (a) rapid heating to above the critical temperature A I, (b) a short time holding at temperature, and (c) cooling below the martensite start ( Ms ) temperature. Some processes also include a short time tempering below 5 00°C after cooling from above Al to improve the ductility and toughness of the steel at the expense of tensile strength. The microstructure of most steels prior to continuous annealing consists of ferrite, pearlite, and grain boundary iron carbides (1, 2, 11-20). Some nominal compositions are listed in Table 1. The rate of heating is far less critical ( 12) than heating temperature, time, or cooling rate. Intercritical heating was preferred by most investi­ gators although steels were also produced by heating supercritically, i.e. above A3 (12, 14). Intercritical heating temperatures offer the inherent control of volume fraction and composition of the ferrite and austenite as dictated by the equilibrium phase diagram (Figure 5). The time at temperature ranged from a few seconds to a few minutes and it is not clear whether equilibrium conditions were attained. The kinetics of austenite formation in these steels have not been investigated but the short heating times suggest that the austenite is probably nonhomoge­ neous and composition gradients may exist. 800 � 600 P!.­ .2 e 400 !. E GI ... 200 Figure 4 Schematic representation of oL-------�-- various steps on the continuous an­ Time- nealing process.
    • 250 RASHID Table 1 Nominal compositions of some dual phase steels Maximum Reference temperature" Cooling Composition, wt.% Continuous Annealed Hayami (I) Air cool 0.09 C, 0.92 Si, 0.97 Mn, 0.32 Cr Matsuoka and Yamomouri (J I) I Water quench 0.07 C, 0.39 Si, 2.96 Mn,by University of Tennessee - Knoxville - Hodges Library on 05/22/12. For personal use only. Rashid (2) I Air cool 0.12 C, 0.51 Si, 1.46 Mn, 0.11 V Bucher (14) I&S Fast air cool 0.11 C, 0.5 Si, 1.4 Mn, 0.06 V Annu. Rev. Mater. Sci. 1981.11:245-266. Downloaded from www.annualreviews.org As-Rolled Coldren (16) S Air cool 0.06 C, 0.9 Si, 1.2 Mn, 0.6 Cr, 0.4 Mo Batch-Annealed Parker (23) Furnace cool (14°Cjhr) 0.08 C, 0.3 Si, 3.0 Mn a 1= intercriticaJ, S= supercriticaJ. Cooling rates in the range between air cooling and water quenching have been used to produce dual phase microstructures. Faster cooling rates are required for steels with lower hardenability. The nonmicroal­ loyed Si-Mn steels are usually produced by water quenching, while slower cooling rates have been used for the microalloyed compositions. Slower cooling rates produce better strength-ductility combinations and are generally preferred. Higher cooling rates induce a larger number of lattice defects and residual stresses into the matrix and may reduce Typical Carbon Content I of Duol Phose Steels I Austenite, Y v I o I i :> 0700 OJ Q. E " I- 600 500 Percent Carbon Figure 5 Schematic representation of a portion of the iron-carbon phase diagram.
    • DUAL PHASE STEELS 251 ductility slightly. However, ductility can be improved by tempering the steel. Some researchers have attempted ( 19, 20) to use existing carbon equivalent formulas to identify compositions and cooling rates that will produce the desired mechanical properties. Such formulae are, at best, only rough indicators of the expected mechanical properties, but they provide a good screening mechanism to test steel compositions and processing parameters. Many unknown parameters exist, however, andby University of Tennessee - Knoxville - Hodges Library on 05/22/12. For personal use only. reliable predictive techniques for steel composition and cooling rates that Annu. Rev. Mater. Sci. 1981.11:245-266. Downloaded from www.annualreviews.org would produce the optimum microstructure have yet to be perfected. As-Rolled In the as-rolled process (16), the steel composition is chosen (Table I) such that 80- 90% of the steel transforms to ferrite after the final roll pass in normal conventional hot rolling and before entering the coiler. The remaining 10-20% does not transform until much later, during slow cooling in the coil. This is possible with steel compositions that exhibit certain special characteristics in their continuous cooling transformation (CCT) diagram (Figure 6): namely (a) an elongated ferrite C-curve, i.e. the ability to form very large amounts of ferrite over a wide range of cooling rates on the run-out table; (b) a suppressed (delayed) pearlite nose and high pearlite finish temperature to ensure avoidance of pearlite formation during cooling to the coiling temperature; and (c) a gap between pearlite and bainitic regions to provide a temperature range within which no transformation occurs, permitting sufficient time for the steel to be coiled. For the composition listed in Table I, the range of 1000 To I AC3 ACI u 800 ° A € 600 Z 2 � 400 Ms E - III M .... 200 coiling window 0 10 0 101 102 103 104 105 Time to Cool from 960°C Figure 6 Continuous cooling transformation diagram of an as-rolled dual phase steel (A = austenite, PF=poJygonal ferrite, P-pearlite, BF=bainitic ferrite, M=martensite of average C content, M=martensite from carbon-enriched austenite, Ta =austenitization temperature) (16).
    • 252 RASHID cooling rates through the ferrite region appears to be rather wide and produces the desired microstructure. Batch Annealed Dual phase steels have also been produced by batch annealing techniques (23) modified for heating in the intercritical temperature range. The very slow cooling rates inherent in this approach (several days to cool to room temperature) necessitate the use of steels with very high alloy contentby University of Tennessee - Knoxville - Hodges Library on 05/22/12. For personal use only. (Mn and Mo) and high hardenability. This approach is presently the least Annu. Rev. Mater. Sci. 1981.11:245-266. Downloaded from www.annualreviews.org researched of the three. TRANSFORMATION MECHANISM Continuous Annealed The mechanisms by which dual phase microstructures are formed are reasonably well understood and can be explained with reference to steel microstructure prior to heating above the critical temperature. As men­ tioned earlier, the starting steel consists of a ferrite matrix with grain boundary iron carbides and small islands of pearlite (Figure 7). Microal­ loyed steels also contain microalloy carbonitrides uniformly distributed throughout the ferrite matrix (Figure 8). Ferrite Pearlite Iron Carbide Pearlite 5 m Figure 7 Scanning electron micrograph of a high strength, low alloy steel. The microstruc­ ture consists of a fine-grained ferrite matrix, grain boundary iron carbides, and islands of pearlite (2).
    • DUAL PHASE STEELS 253by University of Tennessee - Knoxville - Hodges Library on 05/22/12. For personal use only. Annu. Rev. Mater. Sci. 1981.11:245-266. Downloaded from www.annualreviews.org Figure 8 Bright-field micrograph revealing carbonitride precipitates distributed in the . ferrite. Arrows I, 2 and 3 point to precipitates of different sizes. Arrows 2 point to medium-sized carbonitrides precipitated at low angle boundaries, for example, separating regions L and N. Arrows P reveal dislocations pinned by precipitates (26). Upon heating the steel above the critical temperature (Figure 5), islands of carbon-rich, nonequilibrium austenite form at the carbide locations. The heating temperature determines the carbon content and volume fraction of austenite that can exist in equilibrium with ferrite. Given sufficient time at temperature the austenite nuclei grow until this criterion is fulfilled. In steels heated to just above the critical tempera­ ture, the proportion of austenite formed is at a minimum and it has a high carbon content because carbon is more soluble in austenite than in ferrite (Figure 5). At higher annealing temperatures, the volume fraction of austenite is larger and it has a lower carbon content and hence lower hardenability. Steels heated supercritically transform entirely to austenite of the carbon content of the steel. The composition of the austenite is also influenced by other alloying elements in the steel. The presence of Si in the ferrite promotes carbon migration from ferrite to the austenite (24) thereby adding to the carbon content of the austenite, while Mn partitions preferentially to the austenite and increases its hardenability. The austenite composition is usually nonhomogeneous and concentration gradients exist because of the short heating times involved. The substructure and composition of the untrans­ formed ferrite that coexists with the austenite at elevated temperature (2, 26) are also modified. Carbon partitions out of the ferrite, when present microalloy precipitates coarsen, and dislocations rearrange themselves into low energy configurations. The transformation product that forms upon cooling the steel back to room temperature depends on the austenite composition and cooling
    • 254 RASHID rate. Si suppresses the pearlite transformation (25) while C and Mn stabilize the austenite and lower the Ms temperature. At rapid cooling rates all the austenite transforms to martensite. At slower rates, depend­ ing on austenite hardenability, various proportions transform to marten­ site, bainite, and ferrite, with some austenite remaining untransformed (retained austenite). High carbon austenite transforms to twinned martensite while low carbon austenite transforms to lath martensite; the former martensite is stronger than the latter. The volume change andby University of Tennessee - Knoxville - Hodges Library on 05/22/12. For personal use only. shear accompanying the austenite� martensite transformation generates Annu. Rev. Mater. Sci. 1981.11:245-266. Downloaded from www.annualreviews.org numerous new mobile dislocations ( 1, 2) in the surrounding ferrite matrix. The ferrite matrix consists of transformed and untransformed ferrite. As mentioned previously, the untransformed ferrite has low strength and is relatively free of interstitials such as carbon and nitrogen, which have either diffused out of the ferrite or remain in the ferrite and contribute to precipitate coarsening. The transformed ferrite formed when the austenite is cooled below the critical temperature is very similar to the ferrite of the starting steel (26) in dislocation substructure, carbide distribution, and strength. Untransformed ferrite is not expected in supercritically heated steels but very short heating times sometimes preclude complete austenitization and some untransformed ferrite might be observed. As-Rolled Most of the hot rolling is done in the austenite range of the Fe-C phase diagram and the finish rolling temperature is usually in the intercritical region. As mentioned previously, 80-90% of the austenite transforms to ferrite on the run-out table and the remainder transforms in the coil to transformation products similar to those observed after continuous an­ nealing. Some autotempering of the martensite and decomposition of retained austenite may be expected during cooling in the coil. Batch Annealed The transformation mechanisms here are similar to those observed dur­ ing continuous annealing but the grain size and substructure are char­ acteristic of the much slower cooling rates involved. The various phase transformations discussed produce a microstructure consisting of untransformed ferrite, transformed ferrite, martensite, bainite, retained austenite, and carbide precipitates. The relative volume fraction, morphology, distribution, composition, and mechanical prop­ erty of each constituent is governed by steel composition and processing parameters, and determines the deformation behavior of the steel. Defor­ mation of the steel itself can also induce some phase transformations.
    • DUAL PHASE STEELS 255 DEFORMATION BEHAVIOR The deformation behavior of dual phase steels is quite complex. A thorough understanding of the interactions between the various micro­ constituents discussed and their influence on mechanical properties is lacking, but steel deformation behavior can be explained in generalized terms. The deformation behavior of most metals, especially plain carbon steel,by University of Tennessee - Knoxville - Hodges Library on 05/22/12. For personal use only. may be described (27) in terms of a simple empirical relationship between Annu. Rev. Mater. Sci. 1981.11:245-266. Downloaded from www.annualreviews.org the stress-strain data obtained in a tension test, namely a=Ke; 1. or log a=log K+n log€p 2. where a is the true stress, K is the strength constant or a at lOp = 1.0, lOp is the true plastic strain, and n is the strain hardening exponent, a measure of the ability of the metal to distribute strain. The true strain at maximum load, fu also called the true uniform strain, will numerically equal n when Equation I is satisfied (27). When Equation I is satisfied, a plot of log a vs log lOp will be a straight line with slope n. Such behavior is attributed to truly uniform plastic deformation behavior. Equation I is not satisfied for dual phase steels (28). A plot of log (J VS log fp (Figure 9) deviates substantially from linearity, which suggests that 5.2,------ ---, • Ferrite-Pearlite Cooling Rate.o IC 5.1 • 5 09 5.0 -12 014 Figure 9 Variation in plots of log true stress vs. log true plastic strain for steels with the same composition but 4.7 with different microstructures. The ferrite-pearlite steel was heated at 4.6 --_---__=--="__.J..-_-:-o--__ .. .... 788°C for 3 min and cooled to room -3.0 -1.0 temperature at the rates shown to Log True Plastic Strain produce dual phase steels.
    • 256 RASHID several deformation processes are operative in this steel. At least two approaches have been used for detecting changes in deformation behav­ ior. In one approach (28) an incremental value of n, ni is calculated for each segment of the stress-strain curve, with n i being defined as log OJ -log 0j_ . n;(j)=J for /=lto/. 3. og Epu> - 1og Epu_1) •by University of Tennessee - Knoxville - Hodges Library on 05/22/12. For personal use only. Here I is the number of small segments of the stress-strain curve. In the other approach (29) the stress-strain curve is represented by the equation Annu. Rev. Mater. Sci. 1981.11:245-266. Downloaded from www.annualreviews.org 4. where a i s the true stress, 00 i s the true yield stress, Ep is the true plastic strain, and Band m are constants (m is not strain rate sensitivity). Equation 4 can be differentiated (30, 31), logarithms taken on both sides and written as do In- =In Bm+ ( m-l ) ln Ep• 5. d lOp Plots of n i vs lOp (Figure 10) or In dojdEp vs In lOp (Figure 11) are straight lines for ferrite-pearlite steels but delineate several different stages of 0.30 ! en / c: c � 0.20 o :z: c SAE 9 50X / ~ "e .;; g E ; f 0.10 7SA .. OX /// u .5 p"ni 0.2 True Plostic 5troin, Ep Figure 10 Variation of incremental strain-hardening rate, ni with increasing true plastic strain in various steels: ni is relatively constant in the ferrite-pearlite steels but not in the dual phase steel (28).
    • DUAL PHASE STEELS 257 Annealed 4min. @ 810°C �, � , " " 53"e/sec. �" ...- 1//by University of Tennessee - Knoxville - Hodges Library on 05/22/12. For personal use only. 20"C/sec .� Annu. Rev. Mater. Sci. 1981.11:245-266. Downloaded from www.annualreviews.org �Uniform Wt. % &. Strain C 0.15 Mn 1.37 �� Si 0.27 Figure II Plot of Inda/dEp vs In Ep for a N 0.0048 dual phase steel of indicated composition and annealing conditions for different True Plastic Strain cooling rates (29). strain hardening in the deformation behavior of dual phase steels, confirming earlier indications that multistage, nonhomogeneous deforma­ tion occurs in these steels in contrast with homogeneous deformation in ferrite-pearlite steels. Some of these deformation processes can be identified by correlating certain features of the tensile stress-strain curve with the deformation behavior of particular microconstituents in the steel. As mentioned earlier the yield strength in dual phase steels is only 0.5-0.6 of the tensile strength and the yield point elongation is absent. The work-hardening rate is very high at low strains and decreases with increasing strain. The total elongation is higher than that in ferrite-pearlite steels of similar tensile strength. Each of these characteristics can be traced to the deformation behavior of one or more microconstituents described previ­ ously. The shear and volume change accompanying the austenite _ martensite transformation on cooling from above the critical temperature generate numerous free mobile dislocations in the surrounding ferrite matrix. Upon application of a load, the free dislocations move at stresses much lower than required to move restrained dislocations, such as those commonly observed in ferrite-pearlite steels. Hence, dual phase steels yield or commence plastic flow at much lower stresses compared to the ferrite-pearlite steels of equivalent tensile strength. Furthermore, since interstitial solutes such as carbon and nitrogen have either diffused out of the untransformed ferrite or are in combined form, solute-dislocation interactions are severely reduced; an initial threshold stress is not needed to break the dislocation away from the solutes and, hence, no yield point elongation is observed. Upon continued application of tensile load,
    • 258 RASHID plastic flow continues in the microconstituent with the lowest yield strength, this being the untransformed ferrite. After the untransformed ferrite has work-hardened to the yield strength of the transformed ferrite, both phases presumably deform and work-harden simultaneously (2, 26, 28). Metals work-harden because mobile dislocations interact with other dislocations, solutes, precipitates and other microconstituents. This is also true in dual phase steels. But the magnitude of the work-hardeningby University of Tennessee - Knoxville - Hodges Library on 05/22/12. For personal use only. rate observed at low strains is too large to be explained by dislocation Annu. Rev. Mater. Sci. 1981.11:245-266. Downloaded from www.annualreviews.org interactions alone, and could be caused by the added contribution due to deformation of retained austenite. The strain-induced transformation of austenite to martensite increases the ductility and work hardening rate in - several material systems (32, 33). This phenomenon, called transforma­ tion-induced plasticity or TRIP, also occurs in dual phase steels (34, 35). Dual phase steels contain as much as 10 vol% retained austenite, which transforms to martensite in direct rel�tion to increasing strain in a tension test (34). Most of the transformation occurs at low strains, nears completion at about 10% strain and could enhance the work-hardening , rate because of dislocations interacting with other dislocations and with the strain-induced martensite. Deformation obviously continues in the untransformed and trans­ formed ferrites beyond 10% strain, but because of differences in the strengths of the two phases, the work-hardening rate is not constant but continues to decrease (Figure 10) until a constant incremental work­ hardening rate is approached at strains just prior to uniform strain. The deformation characteristics of the ferrite phases are, of course, influenced by composition and process variables. These effects are treated in detail in the next section. Martensite is usually regarded as a nondeformable hard phase that contributes primarily to steel strength. However, martensite deformation has been observed (Figure 12) in dual phase steels at very high strains (28). Concentration gradients in the austenite during continuous annealing evidently produce martensite of varied composition and strength, some of which deforms and contributes differ­ ently to work-hardening behavior than nondeformable martensite. Fur­ thermore, the strain-induced martensite would have yet different defor­ mation characteristics from the original martensite and also contribute to the observed deformation behavior. The higher total elongation of dual phase steels is a consequence of the multistage deformation behavior just described. When dual phase steel is strained, slip leading to deformation occurs first in the constituent with the lowest yield strength. When this constituent work-hardens to the yield strength of the second constituent, plastic flow begins to occur in it. This continues until all constituents are involved in the deformation
    • DUAL PHASE STEELS 259 Fradured Martensite F rriteby University of Tennessee - Knoxville - Hodges Library on 05/22/12. For personal use only. Annu. Rev. Mater. Sci. 1981.11:245-266. Downloaded from www.annualreviews.org Void Martensite Figure J2 Scanning electron micrograph of a dual phase steel at fracture initiation. Both ferrite and martensite grains are elongated. Some martensite islands are fractured and voids are formed at the ferrite-martensite interfaces (28). process and are work-hardened to their maximum extent. Since deforma­ tion is distributed among several constituents, strain is distributed more uniformly and localized necking is delayed, resulting in far better forma­ bility than ferrite-pearlite steels of similar tensile strength. Fracture initiation has also been contrasted in dual phase and ferrite­ pearlite steels. In the latter steels, deformation is restricted to the ferrite phase with no obvious deformation of the iron carbides. When the ferrite is work-hardened to its limit, voids form preferentially at the ferrite-iron carbide interfaces (28) and failure is initiated. In dual phase steels, deformation occurs in the microconstituents in a sequence related to their yield strength. Martensite, being the strongest constituent, does not deform until all other consitituents have deformed and are highly strained. Voids leading to failure form at the ferrite martensite interface but are not nucleated until more extensive deformation has occurred and the martensite is also highly strained (Figure 12). COMPOSITION-STRUCTURE-PROPERTY RELATIONSHIPS Dual phase steels have been modeled as two-phase composites: the ferrite is treated as a homogeneous ductile matrix phase and the "martensite" (or transformation product consisting of martensite, retained austenite,
    • 260 RASHID and bainite) is treated as a high strength reinforcing component with homogeneous mechanical properties. In spite of these simplifying as­ sumptions, the strength of the composite (dual phase steel) has been predicted with reasonable success (36- 38) using the simple rule of mixtures, namely 0c =(1- Vr)om + V.o" where 0c om and Or are the strengths of the composite, the matrix and reinforcement, and v,. is the volume fraction of the reinforcement. The good agreement with experimental data suggests either an insensi­by University of Tennessee - Knoxville - Hodges Library on 05/22/12. For personal use only. tivity in dual phase steels to nonhomogeneous mechanical properties of Annu. Rev. Mater. Sci. 1981.11:245-266. Downloaded from www.annualreviews.org the matrix and reinforcement, a cancellation of errors or just a fortunate coincidence of the assumptions made in the model. Furthermore, these models suggest the rather obvious: optimum properties can be developed in the dual phase steel when the matrix and the reinforcement individu­ ally have maximum strength as well as ductility (38). Since "martensite" is the principal load-bearing constituent, various attempts have been made to correlate volume percent of "martensite" with steel strength. The two are linearly related (Figure 13), independent of "martensite" carbon content (39, 40). Further work showed (41), however, that carbon content is important and that separate linear relationships exist (Figure 14) between yield and tensile strength and percent "martensite" of constant carbon content. A smaller volume of high carbon martensite produces the same strength as a larger volume of martensite of lower carbon content, the high carbon martensite, of course, being stronger than the low carbon martensite. The reason for this apparent contradiction is not clear, but is probably due to scatter in experimental data (Figure 13) and because the "martensite" reported in the earlier work (39, 40) actually contained varying amounts of retained austenite depending on intercritical annealing temperature. The strength could also be affected by variations in the p r ecipitation in the ferrite due to supersaturation. These variations result from the different annealing temperatures used to produce different volumes of martensite. Besides strengthening with carbon or alloy elements, martensite strength can also be increased by decreasing its particle size. This is ensured, in part, by continuously annealing steels with a fine grain size. With a small grain size, the grain boundary iron carbides are proportionately small and this produces correspondingly small martensite islands after trans­ formation. Several types of "martensite" distributions have been ob­ served but the one most conducive to homogeneous deformation is a uniform distribution of very small, disconnected "martensite" islands located at ferrite grain boundary intersections. Sometimes, larger islands are located further apart and some steels are partially banded, with the "martensite" content being higher in the bands. Continuous martensite
    • DUAL PHASE STEELS 261by University of Tennessee - Knoxville - Hodges Library on 05/22/12. For personal use only. Annu. Rev. Mater. Sci. 1981.11:245-266. Downloaded from www.annualreviews.org Martensite, 7. Figure 13 The 0.2% flow stress and tensile strength as a function of volume fraction martensite for Fe-Mn-C alloys (39). distributions along grain boundaries are not conducive to good formabil­ ity and are rarely observed. As discussed previously, a portion of the transformation product is retained austenite, which figures significantly in steel deformation behav­ ior. Essentially two morphologies of retained austenite have been re­ ported (26), submicrometer-size particles uniformly distributed within the ferrite, and larger particles (a few micrometers in diameter) located at prior iron carbide locations closely interspersed with the martensite islands. The latter transform to strain-induced martensite while the former are stable to high strain levels (35). Strengthening by the strain- 7.C in Martensite • 0.64 1000-- • 0.42 .. 0.30 200 1-- � � � ield ..�-- Figure 14 Yield and tensile strengths as function of martensite volume fraction for I I I I °0L----ILo--�2�O�--3�O�--�4�O-J Fe-C-Mn-Si steels with different marten­ Martensite Volume Fraction, 7. site carbon contents (4).
    • 262 RASHID induced martensite adds an extra dimension to the multistage deforma­ tion behavior discussed previously. The martensite in most air-cooled dual phase steels is autotempered and is tougher than untempered martensite. The extent of autotempering is lower in water-quenched dual phase steels and ductility is usually improved in these steels at the expense of strength by a short tempering treatment below 500°C. The microstructure of air-cooled dual phase steels appears quite stable (42) on heating below 200°C. However, onby University of Tennessee - Knoxville - Hodges Library on 05/22/12. For personal use only. heating at higher temperatures (below 500°C) tempering of martensite Annu. Rev. Mater. Sci. 1981.11:245-266. Downloaded from www.annualreviews.org and decomposition of retained austenite is expected. These effects are combined with strain aging by interstitial solutes (43, 44) when strained dual phase steels are heated. The yield strength increase due to strain aging seems to be directly related to steel interstitial content but the change in ultimate tensile strength is influenced by both alloying addition and interstitial content (43). The kinetics of strain aging (44, 45) are slower than that observed in plain carbon steel and the activation energies for strain aging are similar to those reported previously for high strength, low alloy steels (46-48). These steels are also susceptible to dynamic strain aging (49) and exhibit serrated flow when tested at slow rates at elevated temperatures. Mechanical property correlations of steel based on structure or mecha­ nical properties of the ferrite phase have not been reported. The two­ component composite model suggests that ferrite strength and ductility should be maximized simultaneously for optimum strength and ductility in the steel. Ferrite strength can be increased by decreasing grain size and by alloying additions Small grains are produced in HSLA steels by . minor addition of carbide forming elements such as V, Nb, or Ti and by controlled thermomechanical processing. Measurable grain growth does not seem to occur when these steels are continuously annealed and the dual phase steels inherit the small grain size; in fact some grain refine­ ment is observed as a result of the various transformations. A fine starting microstructure, therefore, seems to insure a finer grain size in the dual phase steel. Microalloy additions that enhance grain refinement also precipitation­ strengthen the ferrite and reduce its ductility. These opposing effects must be judiciously balanced to obtain the proper combination of mechanical properties. Heating temperature, steel composition, and cool­ ing rate combine to determine the ratio of untransformed and trans­ formed ferrites. The former has lower yield stress than the latter. The untransformed ferrite is "cleaner" or relatively devoid of interstitial solutes, and large incoherent precipitates are too widely dispersed to contribute to strengthening. The transformed ferrite is much stronger but
    • DUAL PHASE STEELS 263 less ductile (26), the morphology, distribution, and density of the pre­ cipitates being similar to the starting material. Slower cooling rates are conducive to forming precipitates and reducing the interstitial content in both ferrites; they also promote carbon and nitrogen partitioning to the austenite islands, which act as sinks. Slower rates result in better strength-ductility combinations and are usually preferred over quenching after continuous annealing. The role of microalloying additions has often been debated. Microal­by University of Tennessee - Knoxville - Hodges Library on 05/22/12. For personal use only. loys have been used in dual phase steels but some claim they are not Annu. Rev. Mater. Sci. 1981.11:245-266. Downloaded from www.annualreviews.org essential for producing dual phase microstructures with the mechanical properties listed in Figure 3 (50). Vanadium has been added more frequently than niobium or titanium and is thought to fulfill a dual function. Besides playing the principal role of grain refinement, V imparts a higher hardenability to the steel than does Nb or Ti. The latter two precipitate as carbonitrides before and during annealing and there­ fore cannot contribute to hardenability. V has a higher solubility in the matrix and to a large extent remains in solution during annealing, thus contributing to hardenability. Vanadium additions promote martensite formation and produce higher tensile strength than that obtained in equivalently processed vanadium-free steels (51). Ferrite strength can also be increased by substitutional solutes such as P, Mn, and Si, with certain maximum limits of a few percent of each element. Each percent of Mn increases the tensile strength 48 MPa although it is added primarily for increasing steel hardenability. Each percent of Si increases tensile strength 150 MPa or more (47, 51) and these contributions are reported to be additive. In addition to tensile strength, substitutional solutes increase the ferrite work-hardening rate with Si being the most effective of the three. A higher work-hardening rate signifies improved strain distribution and hence formability. The relationship between uniform elongation (or work-hardening rate), and tensile strength (47) is shown in Figure 15 for dual phase steels with various Si and P concentrations. Si additions do indeed improve the ductility of dual phase steel at a given tensile strength level. Addition of P increases strength to a lesser degree than Si and decreases ductility faster than it increases strength. The fatigue behavior of dual phase steels has been studied in as-received and prestrained conditions (52) but correlations have not been reported to volume percent or mechanical properties of the micro-constituents. A comparison of the strain-life curves (Figure 16) shows that at high strains the dual phase steels have better life than SAE 980X, a ferrite-pearlite HSLA steel of equivalent tensile strength, while at low strains the opposite is true. In a notched-fatigue test, which may be more representa-
    • 264 RASHID 30 �� � e .,8 0 20 .�� 01 � 2%Si c Dual 0 ....... Si u:; Phase E 0 :t: 10 HSLA :1" , tee .2%P O.1%P Steels c 15 Comparison of observed uni­by University of Tennessee - Knoxville - Hodges Library on 05/22/12. For personal use only. Figure ::::I form elongation as a function of tensile Annu. Rev. Mater. Sci. 1981.11:245-266. Downloaded from www.annualreviews.org 0 400 strength for both Si- and P-containing Tensile Strength, MPa dual phase steels (47). tive of real-world conditions, the dual phase steel behaved nearly the same as the SAE 980X steel (52). The atmospheric corrosion behavior of dual phase steels is generally similar to that of plain carbon steel. No correlation between composition, strength, or microstructure and corro­ sion behavior has yet been reported. The preceding discussion highlighted present understanding of rela­ tionships between composition, microstructure, and certain mechanical properties. Although some relationships have been explained reasonably well, our understanding of these unique new steels is far from being thorough. FUTURE DEVELOPMENT The development of dual phase steel is unique in many ways. It was developed to satisfy a well-defined need. In just a few years the steel progressed from being a laboratory curiosity to being commercially produced and applied. Technology and application of steel are now N O.lc------ -----, � � .,- Dual-Phase " .2 � 0.01 I E � SAE 950X ] c �L SAE 980X """�::::-;:-- __ VI I ---­ 0.001 Plain Carbon Reversals to Failure, 2Nf Figure 16 Strain-life curves for four sheet steels (52).
    • DUAL PHASE STEELS 265 ahead of theoretical understanding.� The steel has essentially a plain carbon steel composition but innovative processing enables substantial improvement in mechanical properties. The deformation mechanisms are different from those of ferrite-pearlite steels. Most prior research has dealt with the final microstructure of dual phase steels. This trend is now being expanded to study the influence of variations in microstructure prior to continuous annealing. The role of microalloys, that of grain refinement, and the dual role of vanadium willby University of Tennessee - Knoxville - Hodges Library on 05/22/12. For personal use only. probably be clarified. Other trends in developmental work are to in­ Annu. Rev. Mater. Sci. 1981.11:245-266. Downloaded from www.annualreviews.org vestigate tempering and the TRIP phenomenon in dual phase steels. SUMMARY Dual phase steels are characterized by a microstructure consIstIng of 75-85 vol% ferrite with the remainder being a mixture of martensite, bainite, and retained austenite. They have better strength-ductility com­ binations than ferrite-pearlite steels of equivalent tensile strength and are presently being used commercially. Publications and research on these unique steels have increased ex­ ponentially since their inception in the mid 1970s. This paper briefly reviews their historical development and present status of the various approaches that have been used to produce the steels. Also, current understanding of the phase transformations that produce the dual phase microstructure and the deformation mechanisms in these steels are discussed. Present understanding of relationships between composition , microstructure, and tensile properties is briefly reviewed, and some trends in developmental work are presented. Literature Cited 8. Grange, R. A. 1968. US Patent 3,502,514 L Hayami, S., Furukawa, T. 1975. Mi­ 9. Bailey, D. J. 1976. SAE Pre print . croalloying 75, pp. 311 -20. Proc. 760715. Warrendale, Pa: Soc. Auto. Con/., Washington DC. New York: Eng. Union Carbide Corp. 10. Bailey, D. J. 1976. US Patent 2. Rashid. M. S. 1977. SAE Trans. 3,930,907; 1975. 3,928,086 86(2):935-46 11. Matsuoka, T., Yamamori, K. 1975. 3. Garber, S. 1963. Iron Steellnst. London Metall. Trans. A 6:1613-22 Spec. Rep. 79:1-6 12. Rashid, M. S. 1976. SAE Trans. 4. Williams, E. W. 1963. Iron Steel InSf. 85(2):938-49 London Spec. Rep. 79:7-12 13. Rashid, M. S. 1978. US Patent 5. Williams, E. W., Davies, L. K. 1963. 4,129,461 Iron Steel Inst. London Spec. Rep. 14. Bucher, J. H., Hamburg, E. G. 1977. 79:13-20 SAE Trans. 86(1):730-39 6. McFarland, W. H. 1965. Trans. AIME 15. Hamburg, E. G., Cryderman, R. L., 233:2028-35 Butler, J. F. 1977. US Patent 4,033,789 7. Grange, R. A. 1970. Prac. 2nd. Int. 16. Coldren, A. P., Tither, G. 1978. J. Can/. Met. Alloys 2:861-76 Met., 30(4):6-9
    • 266 RASHID 17. Davenport, A. T., ed. 1977. Formable Busch, R. 1967. A SM Trans. Q. 60:252 HSLA and Dual Phase Steels, Conf. 33. Hall, 1. A., Zackay, V. F., Parker, E. Proc. Pittsburgh, Pa: Met. Soc. AIME. R. 1 969. Trans. A SM 62:965-76 1 8. Koo, J. Y., Thomas, G. 1977. Metall. 34. Rigsbee, J. M., VanderAhrend, P. J. Trans. A 8:525-28 1 977. See Ref. 1 7, pp. 56- 86 1 9. Dual Phase and Cold Pressing 35. Rao, B. V. N., Rashid, M. S. 1 98 1 . Vanadium Steels in the Automotive Unpublished data Industry. 1978. Proc. Semin. London: 36. Davies, R. G. 1978. Metall. Trans. A Vanitec 9 : 4 1 - 52 20. Kot, R. A., Morris, J. W. 1979. Struc­ 37. Karlsson, B., Sundstrom, B. O. 1 974. ture and Pro perties o Dual Phase f Mater. Sci. Eng. 16: 1 6 1 -68by University of Tennessee - Knoxville - Hodges Library on 05/22/12. For personal use only. . Steels, Conf. Proc. Met. Soc. AIME. 38. Lagneborg, R. 1978. See Ref. 19, pp. Annu. Rev. Mater. Sci. 1981.11:245-266. Downloaded from www.annualreviews.org 21. Koo, J. Y., Rao, B. V. N., Thomas, G. 43-52 1979. Met. Prog. 1 5:66-70 39. Davies, R. G. 1 978. Metall. Trans. A 22. Rao, B. V. N., Thomas, G. 1 980. 9:67 1 - 79 Metall. Trans. A 1 1 :441-57 40. Butler, J. F., Bucher, 1. H. 1978. See 23. Parker, T. E. 1979. Presented at Ann. Ref. 19, pp. 3- 1 2 Meet. Soc. Auto. Eng., Detroit, Mich. 41. Ramos, L . F . , Matlock, D. K., Krauss, 24. Abramowitz, P., Moll, R. A. 1 970. G. 1979. Metall. Trans. A 10:259-61 Metall. Trans. 1 : 1773-75 42. Rashid, M. S. Unpublished data 25. Heheman, R. F. 1 970. Phase Transfor­ 43. Rashid, M. S. 1 98 1 . Metall. Trans. To mation , pp. 410. Metals Park, Ohio: be published ASM 44. Davies, R. G. 1979. Metall. Trans. A 26. Rao, n. V. N., Sachdev, A. K. 1 98 1 . 0 : 1 549-55 Unpublished data 45. Rashid, M. S. 1976. Metall. Trans. A 27. Dieter, G. E. 1 96 1 . Mechanical Metal­ 7:497-503 lurgy. New York: McGraw Hill. S6 46. Rashid, M. S. 1 975 Metall. Trans. A pp. 6: 1 265-68 28. Rashid, M. S., Cprek, E. R. 1978. For­ 47. Davies, R. G. 1 979. Metall. Trans. A mability Topics -Mettalic Materials, 0: I 1 3- 1 8 ASTM STP 649, pp. 174-90. Phila­ 48. Rashid, M. S. Unpublished data delphia: ASTM 49. Sachdev, A. K. Unpublished data 29. Ramos, L. F., Matlock, D. K., Krauss, 50. Ostrom, P., Lonnberg, B., Lindgren, 1. G., Huppi, G. S. 1 979. Metall. Trans. 1979. Res. Re p. IM-1 404. Stockholm: A 0:259-61 Swed. Inst. Met. Res. 30. Jaoult, B. 1 957. J. Mech . Phys. Solids 51. Repas, P. E. 1 978. See Ref. 1 9, pp. 5 :95- 1 14 1 3 - 22 31. Crussard, C . 1 953. Rev. Metall. Paris 52. Sherman, A. M., Davies, R. G. 1 979. 0:697- 7 0 Metall. Trans. A 10:929- 33 32. Zackay, V. F., Parker, E. R., Fahr, D.,
    • ANNUAL REVIEWS Further Quick links to online content Annual Review of Materials Science Volume 11, 1981 CONTENTS PREFATORY CHAPTER A Perspective on Martensitic Nucleation, G. B. Olson and M. Cohen ExPERIMENTAL AND THEORETICAL METHODS Higher Order Elastic Constants of Solids, Yosio Hiki 51 Weak-Bearn Electron Microscopy, D. J. H. Cockayne 75by University of Tennessee - Knoxville - Hodges Library on 05/22/12. For personal use only. Ellipsometry in Thin Film Analysis, J. B. Theeten and D. E. Aspnes 97 Annu. Rev. Mater. Sci. 1981.11:245-266. Downloaded from www.annualreviews.org Electron-Bearn-Induced Currents in Semiconductors, J. I. Hanoka and R. O. Bell 353 Recent Developments in Lattice Imaging of Materials, R. Sinclair 427 Auger Spectroscopy, D. F. Stein and A. Joshi 485 PREPARATION, PROCESSING, AND STRUCTURAL CHANGES The Microchemistry and Microstructure of Portland Cement, P. L. Pratt and H. M. Jennings 123 Continuous Casting of Steel, Shozo Mizoguchi, Tetsuro Ohashi, and Tsuyoshi Saeki 151 Molecular Bearn Epitaxy of III-V Compounds: Technology and Growth Process, Klaus Ploog 171 Dual Phase Steels, M. S. Rashid 245 Microstructure Fabrication in Electronic Devices, B. E. Deal and P. A. Crossley 321 Formability Maps, J. D. Embury and J. L. Duncan 505 The Selective Permeation of Gases Through Polymers, S. A. Stern and H. L. Frisch 523 PROPERTIES AND PHENOMENA Dislocation Creep, M. M. Myshlyaev 31 Ionic Transport in Amorphous Solid Electrolytes, J. L. Souquet 211 Crystallography at High Pressure, John C. Jamieson 233 Predicting the Fatigue Resistance of Welds, F. V. Lawrence, N.-J. Ho, and P. K. Mazumdar 401 SPECIAL MATERIALS Silicon Nitride Ceramics: Composition, Fabrication Pararneters, and Properties, J. Weiss 381 Properties, Preparation, and Device Applications of Indium Phosphide, K. J. Bachmann 441 STRUCTURE Eutectic Solidification in Ceramic Systems, V. S. Stubican and R. C. Bradt 267 Polymer Alloys, J. W. Barlow and D. R. Paul 299 INDEXES Author Index 551 Subject Index 569 Cumulative Index of Contributing Authors, Volumes 7-11 579 Cumulative Index of Chapter Titles, Volumes 7-11 581